Handbook of Multiphase Polymer Systems

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Handbook of Multiphase Polymer Systems Volume 1

Editors

ABDERRAHIM BOUDENNE, LAURENT IBOS, YVES CANDAU Universit´e Paris-Est, Centre d’Etude et de Recherche en Thermique, Environnement et Syst`emes, Cr´eteil, France

AND SABU THOMAS Centre for Nanoscience and Nanotechnology, Mahatma Gandhi University, Kottayam, Kerala, India

A John Wiley & Sons, Ltd., Publication

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This edition first published 2011  C 2011 John Wiley and Sons Ltd Registered office John Wiley & Sons Ltd, The Atrium, Southern Gate, Chichester, West Sussex, PO19 8SQ, United Kingdom For details of our global editorial offices, for customer services and for information about how to apply for permission to reuse the copyright material in this book please see our website at www.wiley.com. The right of the author to be identified as the author of this work has been asserted in accordance with the Copyright, Designs and Patents Act 1988. All rights reserved. No part of this publication may be reproduced, stored in a retrieval system, or transmitted, in any form or by any means, electronic, mechanical, photocopying, recording or otherwise, except as permitted by the UK Copyright, Designs and Patents Act 1988, without the prior permission of the publisher. Wiley also publishes its books in a variety of electronic formats. Some content that appears in print may not be available in electronic books. Designations used by companies to distinguish their products are often claimed as trademarks. All brand names and product names used in this book are trade names, service marks, trademarks or registered trademarks of their respective owners. The publisher is not associated with any product or vendor mentioned in this book. This publication is designed to provide accurate and authoritative information in regard to the subject matter covered. It is sold on the understanding that the publisher is not engaged in rendering professional services. If professional advice or other expert assistance is required, the services of a competent professional should be sought. The publisher and the author make no representations or warranties with respect to the accuracy or completeness of the contents of this work and specifically disclaim all warranties, including without limitation any implied warranties of fitness for a particular purpose. This work is sold with the understanding that the publisher is not engaged in rendering professional services. The advice and strategies contained herein may not be suitable for every situation. In view of ongoing research, equipment modifications, changes in governmental regulations, and the constant flow of information relating to the use of experimental reagents, equipment, and devices, the reader is urged to review and evaluate the information provided in the package insert or instructions for each chemical, piece of equipment, reagent, or device for, among other things, any changes in the instructions or indication of usage and for added warnings and precautions. The fact that an organization or Website is referred to in this work as a citation and/or a potential source of further information does not mean that the author or the publisher endorses the information the organization or Website may provide or recommendations it may make. Further, readers should be aware that Internet Websites listed in this work may have changed or disappeared between when this work was written and when it is read. No warranty may be created or extended by any promotional statements for this work. Neither the publisher nor the author shall be liable for any damages arising herefrom. Library of Congress Cataloging-in-Publication Data Handbook of multiphase polymer systems / editors, Abderrahim Boudenne ... [et al.]. p. cm. Includes bibliographical references and index. ISBN 978-0-470-71420-1 (cloth) – ISBN 978-1-119-97203-7 (ePDF) – ISBN 978-1-119-97202-0 (oBook) 1. Polymeric composites. I. Boudenne, Abderrahim. TA418.9.C6H3426 2012 547 .7–dc22 2011011524 A catalogue record for this book is available from the British Library. Print ISBN: 9780470714201 ePDF ISBN: 9781119972037 oBook ISBN: 9781119972020 ePub ISBN: 9780470714201 Mobi ISBN: 9781119972891 Typeset in 10/12pt Times by Aptara Inc., New Delhi, India

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List of Contributors

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Foreword

xxiii VOLUME 1

1 Physical, Thermophysical and Interfacial Properties of Multiphase Polymer Systems: State of the Art, New Challenges and Opportunities Sabu Thomas, Abderrahim Boudenne, Laurent Ibos and Yves Candau 1.1 Introduction 1.2 Multiphase Polymer Systems 1.2.1 Polymer Blends 1.2.2 Polymer Composites 1.2.3 Polymer Nanocomposites 1.2.4 Polymer Gels 1.2.5 Interpenetrating Polymer Network System (IPNs) 1.3 A Short Survey of the Literature and Applications 1.4 Book Content 1.4.1 Modeling and Computer Simulation of Multiphase Composites: From Nanoscale to Macroscale Properties 1.4.2 Morphological Investigation Techniques 1.4.3 Macroscopic Physical Characterization 1.4.4 Life Cycling 1.5 Future Outlook, New Challenges and Opportunities References 2 Macro, Micro and Nano Mechanics of Multiphase Polymer Systems Alireza S. Sarvestani and Esmaiel Jabbari 2.1 Introduction 2.2 Unentangled Systems 2.2.1 Microscopic Structure 2.2.2 Macroscopic Properties 2.2.3 Results 2.3 Entangled Systems 2.3.1 Microscopic Structure 2.3.2 Macroscopic Properties 2.3.3 Results 2.4 Conclusion Acknowledgements References

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3 Theory and Simulation of Multiphase Polymer Systems Friederike Schmid 3.1 Introduction 3.2 Basic Concepts of Polymer Theory 3.2.1 Fundamental Properties of Polymer Molecules 3.2.2 Coarse-Graining, Part I 3.2.3 Ideal Chains 3.2.4 Interacting Chains 3.2.5 Chain Dynamics 3.3 Theory of Multiphase Polymer Mixtures 3.3.1 Flory-Huggins Theory 3.3.2 Self-consistent Field Theory 3.3.3 Analytical Theories 3.3.4 An Application: Interfaces in Binary Blends 3.4 Simulations of Multiphase Polymer Systems 3.4.1 Coarse-Graining, Part II 3.4.2 Overview of Structural Models 3.4.3 Overview of Dynamical Models 3.4.4 Applications 3.5 Future Challenges Acknowledgements References 4 Interfaces in Multiphase Polymer Systems Gy¨orgy J. Marosi 4.1 Introduction 4.2 Basic Considerations 4.3 Characteristics of Interfacial Layers 4.3.1 Role of Thermodynamic Factors 4.3.2 Role of Kinetic Factors 4.3.3 Relationship Between Interfacial Structure and Mechanical Response 4.4 Interface Modifications: Types and Aims 4.4.1 Interlayers of Controlled Morphology 4.4.2 Interlayers of Modified Segmental Mobility 4.4.3 Interlayers for Improving the Compatibility of the Phases 4.5 Interlayers of Modified Reactivity 4.6 Responsive Interphases 4.6.1 Non-reversibly Adaptive Interphases 4.6.2 Smart Reversibly Adaptive Interphases 4.7 Methods of Interface Analysis 4.8 Conclusions References 5 Manufacturing of Multiphase Polymeric Systems Soney C. George and Sabu Thomas 5.1 Introduction

31 31 32 32 33 34 36 38 39 39 43 49 55 56 56 58 62 65 70 70 71 81 81 82 83 85 87 88 89 90 92 93 100 101 101 103 105 111 112 123 123

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5.2

5.3

5.4

5.5

5.6

5.7

Manufacturing Techniques of Polymer Blends 5.2.1 Solution Blending 5.2.2 Latex Blending 5.2.3 Freeze-drying 5.2.4 Mechanical Blending 5.2.5 Mechano-chemical Blending 5.2.6 Manufacturing of Polymer Blends Using Supercritical Fluids Manufacturing Techniques of Polymer Composites 5.3.1 Hand Layup Process 5.3.2 Spray Layup Process 5.3.3 Vacuum Bag Molding 5.3.4 Resin Transfer Molding 5.3.5 Pultrusion 5.3.6 Filament Winding Process 5.3.7 Reaction Injection Molding 5.3.8 Rotational Molding Manufacturing Techniques of Nanocomposites 5.4.1 Solution Intercalation 5.4.2 In Situ Intercalative Polymerization Method 5.4.3 Melt Intercalation or Melt Blending Method Manufacturing Techniques of Polymer Gels 5.5.1 Microgels 5.5.2 Aerogels 5.5.3 Xerogels 5.5.4 Nanostructured Gels 5.5.5 Topological Networks 5.5.6 Hydrogels Manufacturing Techniques of Interpenetrating Polymer Networks (IPNs) 5.6.1 Full IPNs 5.6.2 Sequential IPNs 5.6.3 Simultaneous Interpenetrating Networks (SINs) 5.6.4 Latex IPNs 5.6.5 Thermoplastic IPNs 5.6.6 Semi-IPNs 5.6.7 Pseudo-IPNs Conclusion and Future Outlook References

6 Macro, Micro and Nanostructured Morphologies of Multiphase Polymer Systems Han-Xiong Huang 6.1 Introduction 6.1.1 Polymer Blends 6.1.2 Polymer and Its Blend Nanocomposites 6.2 Morphology Development Mechanisms of Multiphase Polymer Systems During Processing 6.2.1 Initial Morphology Development in Polymer Blending 6.2.2 Deformation and Breakup of Droplet

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6.2.3 6.2.4

6.3

6.4

Coalescence of Droplet Intercalated, Exfoliated, and Dispersed Mechanism of Organoclay During Melt Mixing Material-Relevant Factors Affecting the Morphology 6.3.1 Viscosity of Components 6.3.2 Elasticity of Components 6.3.3 Interfacial Tension 6.3.4 Compatibilization 6.3.5 Composition 6.3.6 Nanoparticles Processing-Relevant Factors Affecting the Morphology 6.4.1 Flow Field Types 6.4.2 Chaotic Mixing 6.4.3 Mixing Sequence 6.4.4 Processing Parameters Nomenclature Acknowledgements References

7 Mechanical and Viscoelastic Characterization of Multiphase Polymer Systems Poornima Vijayan P., Siby Varghese and Sabu Thomas 7.1 Introduction 7.2 Polymer Blends 7.2.1 Ultimate Mechanical Properties and Modeling 7.2.2 Dynamic Mechanical Properties 7.2.3 Impact Properties 7.2.4 Nanostructured Polymer Blends 7.3 Interpenetrating Polymer Networks (IPNs) 7.3.1 Modeling of Mechanical Properties of IPNs 7.4 Polymer Gels 7.5 Polymer Composites 7.5.1 Mechanical Properties of Polymer Macrocomposites 7.5.2 Mechanical Properties of Polymer Microcomposites 7.5.3 Mechanical Properties of Polymer Nanocomposites 7.5.4 Mechanical Modeling of Polymer Nanocomposites 7.6 Conclusion, Future Trends and Challenges References 8 Rheology and Viscoelasticity of Multiphase Polymer Systems: Blends and Block Copolymers Jean-Charles Majest´e and Antonio Santamar´ıa 8.1 Introduction 8.2 Morphology of Polymer Blends 8.2.1 Morphology Characterization 8.2.2 Effect of Rheological Parameters on Morphology

174 178 183 183 187 192 197 208 212 218 218 221 227 230 231 234 235 251 251 253 253 271 276 279 282 287 290 293 295 298 298 304 307 307

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8.3

8.4

8.5 8.6

8.7

8.8

8.9

Microrheology of Droplet Deformation 8.3.1 Breakup 8.3.2 Coalescence Rheology of Polymer Blends 8.4.1 Specificity of Blend Rheology 8.4.2 Blending Laws and Viscoelasticity Models 8.4.3 Low Frequency Viscoelastic Behavior of Polymer Blends Microphase Separated Block Copolymers 8.5.1 Ordered State and Morphologies in Block Copolymers: The Case of SEBS Triblock Dynamic Viscoelastic Results of SEBS Copolymers 8.6.1 Low and Intermediate Frequency Viscoelastic Behavior 8.6.2 Thermorheological Complexity 8.6.3 Specific Mechanical Relaxation at Low Frequencies Flow-induced Morphological Changes 8.7.1 Order–order Transition and Flow Alignment in Block Copolymers 8.7.2 Flow Alignment in a SEBS Copolymer Capillary Extrusion Rheometry Results of Block Copolymers 8.8.1 General Results of Styrenic Block Copolymers 8.8.2 Viscosity and Flow Instabilities in SEBS Copolymers Summary References

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318 318 321 322 322 326 335 339 339 340 340 341 343 345 345 345 347 347 350 354 354

9 Thermal Analysis of Multiphase Polymer Systems Gy¨orgy J. Marosi, Alfr´ed Menyh´ard, G´eza Regdon Jr. and J´ozsef Varga 9.1 Introduction 9.2 Thermo-optical Microscopy 9.3 Differential Scanning Calorimetry 9.4 Temperature Modulated Differential Scanning Calorimetry 9.5 Micro- and Nanothermal Analysis 9.6 Thermal Gravimetric Analysis and Evolved Gas Analysis 9.7 Conclusions References

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Thermophysical Properties of Multiphase Polymer Systems Abderrahim Boudenne, Laurent Ibos and Yves Candau 10.1 Introduction 10.2 Thermophysical Properties: Short Definitions 10.3 Measurement Techniques 10.3.1 Methods for the Measurement of One Property 10.3.2 Methods for the Simultaneous Measurement of Several Parameters 10.4 Thermophysical Properties of Polymers and Composite Systems 10.4.1 Neat Polymers (Unfilled Systems) 10.4.2 Thermophysical Behavior of Composites 10.5 Summary References

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Electrically Conductive Polymeric Composites and Nanocomposites ˇ Igor Krupa, Jan Prokeˇs, Ivo Kˇrivka and Zdeno Spitalsk´ y 11.1 Introduction 11.2 Theory 11.2.1 Percolation Models 11.3 Electrically Conductive Fillers 11.3.1 Carbon Black 11.3.2 Metallic Fillers 11.3.3 Metallized Fillers 11.3.4 Graphite 11.3.5 Carbon Nanotubes (CNT) 11.3.6 Conducting Polymers 11.3.7 Fillers Coated by Conducting Polymers 11.4 Effect of Processing Conditions on the Electrical Behavior of Composites 11.4.1 Blending of Polymeric Composites 11.4.2 Effects of the Secondary Processing Steps on Conductivity 11.4.3 Effect of Polymer Characteristics on the Electrical Conductivity of Composites 11.4.4 Crystallinity Effect 11.4.5 Effect of Polymer–Filler Interaction 11.4.6 Multiphase Morphology of Polymers and Its Influence on the Conductivity of Composites: Multipercolation Effect 11.5 Applications 11.5.1 EMI 11.5.2 ESD 11.5.3 Electrically Conductive Adhesives 11.5.4 Conductive Rubbers 11.5.5 Semi-Conductive Cable Compounds 11.5.6 Fuel Cells 11.6 Resistance Measurements 11.6.1 Two-Probes Method 11.6.2 Four-Probes Method 11.6.3 Van der Pauw Method 11.6.4 Spreading Resistance of the Contacts 11.6.5 Contact Resistance References

425 425 426 427 432 432 435 438 438 443 449 451 452 452 453 454 454 455 456 457 457 457 458 458 458 458 458 459 460 465 468 470 472

VOLUME 2 12

Dielectric Spectroscopy and Thermally Stimulated Depolarization Current Analysis of Multiphase Polymer Systems Polycarpos Pissis, Apostolos Kyritsis and Daniel Fragiadakis 12.1 Introduction 12.2 Dielectric Techniques 12.2.1 Introduction 12.2.2 Broadband Dielectric Spectroscopy (DS) 12.2.3 Thermally Stimulated Depolarization Current (TSDC) Techniques

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12.3

12.4

12.5

12.6

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Copolymers and Interpenetrating Polymer Networks Based on Poly(alkyl acrylate)s and Poly(alkyl methacrylate)s (Mixing and Phase Separation) 12.3.1 Introduction 12.3.2 Poly(butyl acrylate)-Poly(butyl methacrylate) Sequential Interpenetrating Polymer Networks 12.3.3 Poly(butyl acrylate)-Poly(methyl methacrylate) Interpenetrating Polymer Networks and Copolymers 12.3.4 Poly(ethyl methacrylate)-Poly(hydroxyethyl acrylate) Copolymers 12.3.5 Concluding Remarks Rubber/Silica Nanocomposites (Interfacial Phenomena) 12.4.1 Introduction 12.4.2 TSDC Studies 12.4.3 Broadband DS Studies 12.4.4 Concluding Remarks Polymer Nanocomposites with Conductive Carbon Inclusions (Percolation Phenomena) 12.5.1 Introduction 12.5.2 Analysis of DS Data in Terms of the Dielectric Function 12.5.3 Analysis of DS Data in Terms of ac Conductivity 12.5.4 Concluding Remarks Conclusion Acknowledgements References

Solid-State NMR Spectroscopy of Multiphase Polymer Systems Antonio Mart´ınez-Richa and Regan L. Silvestri 13.1 Introduction to NMR 13.2 Phases in Polymers: Polymer Conformation 13.3 High Resolution 13 C NMR Spectroscopy of Solid Polymers 13.3.1 Chemical Shift 13.3.2 Polyolefines 13.3.3 Polyesters 13.3.4 Carbohydrates 13.3.5 Conducting Polymers 13.3.6 Polymer Blends 13.3.7 Interactions Between Polymers and Low-molecular Weight Compounds 13.3.8 Miscellaneous Polymers 13.4 Additional Nuclei 13.5 NMR Relaxation 13.5.1 NMR Relaxation in the Study of Polymer Blends 13.5.2 NMR Relaxation in the Study of Copolymers 13.5.3 NMR Relaxation in the Study of Polymer Composites 13.5.4 NMR Relaxation in the Study of Polymers for Drug Delivery 13.6 Spin Diffusion 13.7 Concluding Remarks References

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Contents

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ESR Spectroscopy of Multiphase Polymer Systems Sre´cko Vali´c, Mladen Andreis and Damir Klepac 14.1 Introduction 14.2 Theoretical Background 14.3 Copolymers 14.4 Grafted Polymers 14.5 Blends 14.6 Crosslinked Polymers 14.7 Semi-Interpenetrating Networks (SIPNs) 14.8 Composites 14.9 Nanocomposites 14.10 Other Polymer Multiphase Systems 14.11 Conclusion References

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XPS Studies of Multiphase Polymer Systems Mohamed M. Chehimi, Fatma Djouani and Karim Benzarti 15.1 Introduction 15.2 Basic Principles of X-ray Photoelectron Spectroscopy 15.2.1 Photoionization 15.2.2 Surface Specificity of XPS 15.2.3 Spectral Examination and Analysis 15.2.4 Quantification 15.2.5 Determination of Overlayer Thickness 15.2.6 Instrumentation 15.3 Applications of XPS to Polymeric Materials 15.3.1 Polymer Grafts 15.3.2 Colloidal Particles 15.3.3 Epoxy Adhesives 15.3.4 Conductive Polymers 15.3.5 Polymer Blends 15.3.6 Composites 15.3.7 Interpenetrating Polymer Networks 15.3.8 Random and Block Copolymers 15.4 Conclusion Glossary References Light Scattering Studies of Multiphase Polymer Systems Yajiang Huang, Xia Liao, Qi Yang and Guangxian Li 16.1 Introduction 16.2 Light Scattering Technique 16.2.1 Scattering from Multiphase Polymer Systems 16.2.2 Experiment 16.2.3 Intensity Calibration

551 551 555 560 569 569 571 571 573 575 576 578 579 585 585 586 586 587 587 592 592 597 599 600 605 609 613 619 624 627 628 629 630 631 639 639 640 640 642 645

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16.4 16.5

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Phase Behavior of Multiphase Polymer Systems Studied by SALS 16.3.1 Thermodynamics 16.3.2 Phase Separation Dynamics 16.3.3 Reaction-induced Phase Separation 16.3.4 Phase Behavior of Polymer Blends Under Shear Flow 16.3.5 Multi-scale Approaches in Studying the Phase Behavior of Polymer Blends On-line Morphological Characterization of Polymer Blends Light Scattering Characterization of Other Multiphase Polymer Systems 16.5.1 Gelation 16.5.2 Crystallization References

X-ray Scattering Studies on Multiphasic Polymer Systems Z. Z. Denchev and J. C Viana 17.1 Introduction 17.2 Theoretical Background 17.2.1 Microfibrillar Reinforced Composites (MCF): Definition and Preparation 17.2.2 Clay-containing Polymer Nanocomposites 17.2.3 The use of WAXS and SAXS in Characterization of Polymers 17.3 Studies on Multiphase Polymer Systems 17.3.1 Polyamide 6/montmorillonite Nanocomposites 17.3.2 Microfibrillar Composites (MFC) 17.3.3 Immiscible Polymer Blends 17.3.4 Non-conventional Molding of PP Nanocomposites 17.3.5 Stretching of Nanoclay PET Nanocomposite 17.4 Concluding Remarks Acknowledgements References Characterization of Multiphase Polymer Systems by Neutron Scattering Max Wolff 18.1 Introduction 18.2 Method of Neutron Scattering 18.2.1 Scattering Experiment 18.2.2 Born Approximation 18.2.3 Elastic and Quasielastic Scattering 18.2.4 Scattering at Small Momentum Transfer 18.3 Experimental Techniques 18.3.1 Production and Detection of Neutrons 18.3.2 Instrumentation 18.3.3 Grazing Incidence Small Angle Scattering 18.3.4 Comparison of SANS and GISANS for Crystalline Systems 18.4 Recent Experimental Results 18.4.1 Polymer Dynamics 18.4.2 Contrast Variation 18.4.3 Effect of Shear 18.4.4 Near Surface Crystallization of Micelles

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Conclusion Acknowledgements References

Gas Diffusion in Multiphase Polymer Systems Eliane Espuche 19.1 Introduction 19.2 Gas Transport Mechanisms in Dense Polymer Films: Definition of the Transport Parameters 19.3 Multiphase Polymer Systems for Improved Barrier Properties 19.3.1 Introduction 19.3.2 Dispersion of Impermeable Spheres Within a Polymer Matrix 19.3.3 Influence of the Shape of the Dispersed Impermeable Phase: Interest of Oriented Polymer Blends and of the Nanocomposite Approach 19.3.4 Nanocomposites Based on Lamellar Nanofillers 19.3.5 Multilayers 19.3.6 Active Films 19.3.7 Comparison of the Different Ways Used to Improve Barrier Properties 19.4 Multiphase Polymer-based Systems for Improved Selectivity 19.4.1 Introduction 19.4.2 Organic–inorganic Materials for Gas Separation Membranes 19.5 Conclusion References Nondestructive Testing of Composite Materials Zhongyi Zhang and Mel Richardson 20.1 Introduction 20.2 Failure Mechanisms in Polymer Composites 20.2.1 Matrix Deformation 20.2.2 Fiber–matrix Debonding 20.2.3 Matrix Cracking 20.2.4 Delamination 20.2.5 Fiber Breakage 20.2.6 Combination of Different Failure Modes 20.3 Visual Inspection 20.4 Acoustic Emission 20.5 Ultrasonic Scanning 20.6 Radiography 20.7 Thermography 20.8 Laser Interferometry 20.9 Electronic Shearography 20.10 Optical Deformation and Strain Measurement System 20.11 Summary References

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Ageing and Degradation of Multiphase Polymer Systems Xavier Colin, Gilbert Teyssedre and Magali Fois 21.1 Introduction 21.1.1 Issues Associated With Material Ageing 21.1.2 Classification of Ageing Types 21.2 Physical Ageing 21.2.1 Ageing Induced by Structural Reorganization 21.2.2 Ageing Induced by Solvent Absorption 21.2.3 Ageing Induced by Additive Migration 21.3 Chemical Ageing 21.3.1 General Aspects 21.3.2 Mechanistic Schemes 21.4 Impact of Multiphase Structure on Ageing Processes 21.4.1 Structural Reorganization 21.4.2 Diffusion Controlled Processes 21.5 Practical Impact of Physical Ageing on Use Properties 21.5.1 Water-Induced Mechanical Damages in Composites 21.5.2 Ageing of Electrical Insulations 21.6 Concluding Remarks References Fire Retardancy of Multiphase Polymer Systems Michel Ferriol, Fouad Laoutid and Jos´e-Marie Lopez Cuesta 22.1 Introduction 22.2 Combustion and Flame Retardancy of Polymers 22.2.1 Combustion of Polymers 22.2.2 Flame Retardancy 22.3 Laboratory Fire Testing 22.3.1 Limiting Oxygen Index (LOI) 22.3.2 Epiradiator or ‘Drop Test’ 22.3.3 UL 94 22.3.4 Cone Calorimeter 22.4 Flame Retardant Additives 22.4.1 Hydrated Fillers 22.4.2 Halogenated Flame Retardants 22.4.3 Phosphorus-based Flame Retardants 22.4.4 Nanometric Particles 22.5 Synergistic Effects of Fillers with Flame Retardant Additives 22.5.1 Definition of Synergistic Effects in Flame Retardant Systems 22.5.2 Micronic Fillers and Flame Retardants 22.5.3 Nanometric Fillers and Flame Retardants 22.6 Conclusion Acknowledgements References

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797 797 797 798 799 800 802 803 806 806 816 829 829 831 832 832 833 839 840 843 843 844 844 845 846 846 847 848 848 849 849 850 850 852 854 854 854 856 862 862 863

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Applications of Selected Multiphase Systems Igor Nov´ak, Volkan Cecen and Vladim´ır Poll´ak 23.1 Introduction 23.2 Construction Applications 23.2.1 Automotive Applications 23.2.2 Marine Applications 23.2.3 Other Applications 23.3 Aeronautics and Spacecraft Applications 23.3.1 Aeronautics Applications 23.3.2 Spacecraft Applications 23.4 Human Medicine Applications 23.4.1 Musculoskeletal and Bone Applications 23.4.2 Dentistry Applications 23.5 Electrical and Electronic Applications 23.6 Conclusion References

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Waste Management, Recycling and Regeneration of Filled Polymers Jos´e-Marie Lopez Cuesta, Didier Perrin, and Rodolphe Sonnier 24.1 Introduction 24.2 Identification and Sorting 24.3 Separation of Components 24.3.1 Mechanical Separation 24.3.2 Dissolution of Resin 24.4 Feedstock Recycling 24.5 Thermal Processes 24.6 Mechanical Recycling of Filled Thermoplastics 24.6.1 Degradation During Reprocessing of Filled Thermoplastics and Influence of Interfacial Agents 24.6.2 Properties of Recycled Filled or Reinforced Thermoplastics 24.6.3 Recycling of Polymer Nanocomposites 24.7 Waste Management of Glass Fiber-reinforced Thermoset Plastics 24.7.1 Waste Management and International Context 24.7.2 Feedstock Recycling by Pyrolysis 24.7.3 Solvolysis or Chemical Recycling 24.7.4 Mechanical Recycling of Glass-reinforced Thermoset Composites 24.8 Conclusion References

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Nanoparticle Reinforcement of Elastomers and Some Other Types of Polymers James E. Mark 25.1 Introduction 25.2 Fillers in Elastomers 25.2.1 Generation of Approximately Spherical Particles 25.2.2 Glassy Particles Deformable into Ellipsoidal Shapes 25.2.3 Layered Fillers

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25.3 25.4 25.5 25.6 25.7 25.8 25.9

Index

25.2.4 Magnetic Particles 25.2.5 Polyhedral Oligomeric Silsesquioxanes (POSS) 25.2.6 Nanotubes 25.2.7 Dual Fillers 25.2.8 Porous Fillers 25.2.9 Fillers with Controlled Interfaces 25.2.10 Silicification and Biosilicification 25.2.11 Theory and Simulations on Filler Reinforcement Nanoparticles in Glassy Polymers Nanoparticles in Partially-Crystalline Polymers Nanoparticles in Naturally-Occurring Polymers Nanoparticles in Relatively-Rigid Polymers Nanoparticles in Thermoset Polymers Conclusions Acknowledgements References

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965 965 966 966 967 967 967 968 970 970 971 972 972 973 973 973 981

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Mladen Andreis, Rudjer Boˇskovi´c Institute, Zagreb, Croatia Karim Benzarti, Laboratoire Central des Ponts et Chaussees, Paris, France Abderrahim Boudenne, Universit´e Paris-Est, CERTES EA 3481 – Centre d’Etude et de Recherche en Thermique, Environnement et Syst`emes, Cr´eteil, France Yves Candau, Universit´e Paris-Est, CERTES EA 3481 – Centre d’Etude et de Recherche en Thermique, Environnement et Syst`emes, Cr´eteil, France Volkan Cecen, Department of Mechanical Engineering, Dokuz Eylul University, Bornova, Izmir, Turkey Mohamed M. Chehim, ITODYS, Univeresity Paris Diderot and CNRS, Paris, France Xavier Colin, PIMM, Arts et M´etiers Paris Tech, Paris, France Jos´e-Marie Lopez Cuesta, CMGD, Ecole des Mines d’Ales, Ales, France Z.Z. Denchev, Institute for Polymers and Composites, University of Minho, Minho, Portugal Fatma Djouan, ITODYS, University Paris Diderot and CNRS, Paris, France Elian Espuche, Ing´enierie des Mat´eriaux Polym`eres, UMR CNRS 5223, IMP@UCB, Universit´e de Lyon, Universit´e Lyon 1, France Michel Ferriol, LMOPS, Universit´ePaul Verlaine Metz, Sain-Avold, France Magali Fois, Universit´e Paris-Est, CERTES EA 3481 – Centre d’Etude et de Recherche en Thermique, Environnement et Syst`emes, Cr´eteil, France Daniel Fragiadakis, Naval Research Laboratory, Washington, DC, USA Soney C. George, Department of Basic Science, Amal Jyothi College of Engineering, Kerala, India Han-Xiong Huang, Laboratory for Micro Molding and Polymer Rheology, South China University of Technology, Guangzhou, China Yajiang Huang, College of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Sichuan, China Esmaiel Jabbari, Department of Chemical Engineering, University of South Carolina, Columbia, USA Laurent Ibos, Universit´e Paris-Est, CERTES EA 3481 – Centre d’Etude et de Recherche en Thermique, Environnement et Syst`emes, Cr´eteil, France Damir Klepac, School of Medicine, University of Rijeka, Rijeka, Croatia Ivo Kˇrivka, Department of Macromolecular Physics, Faculty of Mathematics and Physics, Charles University in Prague, Prague, Czech Republic

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Igor Krupa, Polymer Institute, Slovak Academy of Sciences, D´ubravsk´a, Bratislava, Slovakia Apostolos Kyritsis, National Technical University of Athens, Athens, Greece Fouad Laoutid, Materia Nova Asbl, Mons, Belgium Guangxian Li, College of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Sichuan, China Xia Liao, College of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Sichuan, China Jean-Charles Majest´e, Laboratoire de Rh´eologie des Mati`eres Plastiques, CNRS, St Etienne, France James E. Mark, Department of Chemistry and the Polymer Research Center, The University of Cincinnati, Cincinnati, Ohio, USA Gy¨orgy J. Marosi, Budapest University of Technology and Economics, Budapest, Hungary Antonio Mart´ınez-Richa, Departamento de Quimica, Universidad de Guanajuato, Guanajuato, Mexico Alfr´ed Menyhard, Jr., Budapest University of Technology and Economics, Budapest, Hungary Igor Nov´ak, Polymer Institute, Slovak Academy of Sciences, Bratislava, Slovakia Didier Perrin, CMGD, Ecole des Mines d’Ales, Ales, France Polycarpos Pissis, National Technical University of Athens, Athens, Greece Vladimir Poll´ak, Polymer Institute, Slovak Academy of Sciences, Bratislava, Slovakia Jan Prokeˇs, Department of Macromolecular Physics, Faculty of Mathematics and Physics, Charles University in Prague, Prague, Czech Republic G´eza Regdon, Jr., University of Szeged, Szeged, Hungary Mel Richardson, Department of Mechanical and Design Engineering, University of Portsmouth, Portsmouth, UK Antonio Santamar´ıa, Polymer Science and Technology Department, Faculty of Chemistry, University of the Basque Country, San Sebasti´an, Spain Alireza S. Sarvestani, Department of Mechanical Engineering, University of Maine, Orono, Maine, USA Friederike Schmid, Institute of Physics, Johannes-Gutenberg Universit¨at Mainz, Germany Regan L. Silvestri, Department of Chemistry, Baldwin-Wallace College, Berea, Ohio, USA Rodolphe Sonnier, CMGD, Ecole des Mines d’Ales, Ales, France ˇ Zdeno Spitalsk´ y, Polymer Institute, Slovak Academy of Sciences, D´ubravsk´a, Bratislava, Slovakia Gilbert Teyssedre, Laplace Universit´e Paul Sabatier, Toulouse, France Sabu Thomas, Centre for Nanoscience and Nanotechnology, Mahatma Gandhi University, Kerala, India Sre´cko Vali´c, University of Rijeka, Rijeka, Croatia, and Rudjer Boˇskovi´c Institute, Zagreb, Croatia J´ozsef Varga, Budapest University of Technology and Economics, Budapest, Hungary

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Siby Varghese, Rubber Research Institute of India, Kottayam, Kerala, India J.C. Viana, Institute for Polymers and Composites, University of Minho, Minho, Portugal Poornima Vijayan P, School of Chemical Sciences, Mahatma Gandhi University, Kerala, India Max Wolff, Department of Phyics, Uppsala University, Uppsala, Sweden Qi Yang, College of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Sichuan, China Zhongyi Zhang, Advanced Polymer and Composites (APC) Research Group, Department of Mechanical and Design Engineering, University of Portsmouth, Portsmouth, UK

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Multiphase polymer systems have been the focus of recent research and have become an important issue from both the industrial and fundamental points of view. The scientific literature devoted to multiphase polymer systems is large and growing as it covers a wide range of materials such as composites, blends, alloys, gels and Interpenetrating Polymer Networks. During the last two decades, major opportunities have appeared due to the possibility of tuning the different relevant length scales with the promise to produce a new generation of materials displaying enhanced physical, mechanical, thermal, electrical, magnetic, and optical properties. In spite of these intensive investigations, there are still many unresolved problems in this field. One of the main issues is the influence of the shape, size and dispersion of the particles in the polymer matrix on the macroscopic behavior of the resulting material. There are many factors which control the dispersion, and one of them is the interaction between the particles and the polymer phase. Describing the interactions between the various components, the physical attributes of polymers and particles, the physical, thermophysical and interfacial properties in a comprehensive universal scheme remains a challenge. This approach requires collecting a large number of experimental data that can be obtained only by using various and complementary experimental techniques. Investigations in this field cover different topics, such as polymer blends and composites and nanocomposites reinforcement, barrier properties, flame resistance, electro-optical properties, etc. Part of these multiphase polymer materials belong to the so-called smart materials which are materials that have one or more properties that can be significantly changed in a controlled fashion by external stimuli. The key to the success of these smart materials hinges on the ability to exploit the potential of nano-structuring in the final product. This book discusses many of the recent advances that have been made in the field of morphological, interfacial, physical, rheological and thermophysical properties of multiphase polymer systems. Its content is original in the sense that it pays particular attention to the different length scales (macro, micro and nano) which are necessary for a full understanding of the structure–property relationships of multiphase polymer systems. It gives a good survey of the manufacturing and processing techniques needed to produce these materials. A complete state-of-the-art is given of all the currently available techniques for the characterization of these multiphase systems over a wide range of time and space scales. Theoretical prediction of the properties of multiphase polymer systems is also very important, not only to analyze and optimize material performance, but also to design new material. This book gives a critical summary of the existing major analytical and numerical approaches dealing with material property modeling. Most of the applications of these smart materials are also reviewed which shows clearly their important impact on a wide range of the new technologies which are currently used in our daily life. Finally, the ageing, degradation and recycling of multiphase polymer systems is not forgotten and some routes are proposed to avoid environmental contamination. The 52 contributors of this book are all leading researchers in their respective fields, and I warmly congratulate the editors Abderrahim Boudenne, Yves Candau, Laurent Ibos and Sabu Thomas for bringing them together to produce this original and important book dealing on multiphase polymer systems.

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I am quite convinced that this book will serve as a reference and guide for those who work in this area or wish to learn about these promising new materials.

Dominique Durand Laboratoire de Physicochimie Macromol´eculaire, Equipe de Recherche Associ´ee au Centre National de la Recherche Scientifique, Facult´e des Sciences, Route de Laval, le Mans, France

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1 Physical, Thermophysical and Interfacial Properties of Multiphase Polymer Systems: State of the Art, New Challenges and Opportunities Sabu Thomas Centre for Nanoscience and Nanotechnology, Mahatma Gandhi University, Kerala, India

Abderrahim Boudenne, Laurent Ibos and Yves Candau Universite Paris-Est, Cr´eteil Centre d’Etude et de Recherche en Thermique, Environnement et Systemes, 61 Av. du G´en´eral de Gaulle 94010 Cr´eteil Cedex, France

1.1 Introduction Multicomponent polymer systems find a wide range of applications in each and every phase of our dayto-day life. Continued research has resulted in the development of super performing macro-, micro- and nanostructured polymeric materials. The new emerging fields of micro- and nano-composites have put forward many challenging opportunities for the use of these smart materials. Polymer physicists, chemists, engineers and technologists show great interest in new strategies for developing high-performance multicomponent systems. Recently, polymer nanostructured multiphase systems have gained much interest due to their unique properties. Characterization of the interphase, physical properties and thermophysical properties are crucial for the understanding of the behavior of these smart materials. A comprehensive understanding of these materials is vital for the industrial use of these materials. The main objective of this book is to present a survey of recent advances in the area of multiphase polymer systems covering physical, interfacial and thermophysical properties of these materials. After a short presentation of the different existing multiphase polymer systems, followed by a survey of actual scientific production and of application fields for these materials, we present some of the recent developments in the Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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area of multicomponent polymer systems that will be highlighted all through the book. The chapter ends with a summary of unresolved issues, perspectives and new challenges for the future.

1.2

Multiphase Polymer Systems

Multiphase polymer systems are characterized by the simultaneous presence of several phases, the two-phase system being the simplest case. Many of the materials described by the term multiphase are two-phase systems that may show a multitude of finely dispersed phase domains. The term ‘two-component’ is sometimes used to describe flows in which the phases consist of different chemical substances. Multiphase polymer systems in general include polymer blends, composites, nanocomposites, interpenetrating polymer networks (IPNs), and polymer gels. 1.2.1

Polymer Blends

Polymer blends can be considered as a macroscopically homogeneous mixture of two or more polymeric species with synergistic properties. In most cases, blends are homogenous on scales larger than several times the wavelength of visible light. Blends may be either compatible or incompatible. Polymer blends can be broadly divided into three categories: miscible, partially miscible and immiscible blends. A miscible polymer blend is capable of forming a single phase over certain ranges of temperature, pressure, and composition; also it can be thermodynamically stable or metastable, exhibits a single Tg or optical clarity. An immiscible polymer blend means a multiphase system. Although polymer blending is a very attractive way to obtain new materials, most polymers are immiscible and/or incompatible. Reasons for incompatibility are high interfacial tension and poor interfacial adhesion. In general, a miscible blend of two polymers is going to have properties somewhere between those of the two unblended polymers. Whether or not a single phase exists depends on the chemical structure, molar mass distribution and molecular architecture of the components present. The single phase in a mixture may be confirmed by light scattering, X-ray scattering and neutron scattering. Typical dispersed phase morphology of polymer blend is given in Figure 1.1 [1]. 1.2.2

Polymer Composites

Generally a composite is defined as a multi-component material comprising multiple different (nongaseous) phase domains in which at least one type of phase domain is a continuous phase. In polymer composites, at least one component is a polymer. Fillers such as fibers, particulate fillers, wood fibers, glass fibers and minerals are used as reinforcements in polymeric matrices. 1.2.3

Polymer Nanocomposites

A nanocomposite is a composite in which at least one of the phases has at least one dimension of the order of nanometers, or structures having nano-scale repeat distances between the different phases that make up the material. Polymeric nanocomposites prepared from high aspect ratio fillers such as carbon nanotubes, layered graphite nanofillers etc. achieve significant improvements in mechanical and electrical properties at low filler concentrations, compared to conventional composites [2, 3], without a significant increase in density (see Figure 1.2). The polymer nanocomposites could be prepared by solution mixing process, in situ intercalation process, latex compounding techniques, and melt mixing techniques (see Figure 1.3). In the case of layered clay nanocomposites, one can achieve different types of morphologies as shown in Figure 1.4.

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Physical, Thermophysical and Interfacial Properties of Multiphase Polymer Systems Drops

Double emulsion

(toughness, 1 μm surface modification) Fibers

(toughness and stiffness)

Cocontinuous

(strength, thermal expansion)

3

Laminar

(barrier)

Ordered microphases

(high flow, electrical conductivity, toughness, stiffness)

10 nm

Figure 1.1 Dispersed phase morphology of polymer blends. Reprinted from [1]. Copyright (2000) with permission from John Wiley and Sons.

1.2.4

Polymer Gels

Polymer gels consist of a crosslinked polymer network inflated with a solvent such as water. They have the ability to reversibly swell or shrink (up to 1000 times in volume) due to small changes in their environment (pH, temperature, electric field). The swelling behavior of gels is presented in Figure 1.5. 1.2.5

Interpenetrating Polymer Network System (IPNs)

An interpenetrating polymer network (IPN) is a polymer comprising two or more networks which are at least partially interlaced on a polymer scale but not covalently bonded to each other. The network cannot be separated unless chemical bonds are broken. The two or more networks can be envisioned to be entangled in such a way that they are concatenated and cannot be pulled apart, but not bonded to each other by any chemical bond. There are semi-interpenetrating polymer networks and pseudo-interpenetrating polymer

Figure 1.2

Typical morphology of nanocomposites.

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Synthesis Approaches Melt Intercalation: Co-extrusion Functionalization of the NPs*

SWNT Rope

Tailoring the modifier to the polymer promotes favorable interactions SWNT

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O

n

Organic Modifier

In-situ polymerization with pristine** or functionalized NPs Δ or h ν

Surfactant assisted dispersion of NPs***

Figure 1.3 Synthesis approach to polymer nanocomposites.

networks. A polymer comprising one or more polymer network(s) and one or more linear or branched polymer(s) is characterized by the penetration on a molecular scale of at least one of the networks by at least some of the linear or branched chains. Semi-interpenetrating polymer networks (SIPNs) may be further described by the process by which they are synthesized. These include sequential SIPNs, simultaneous SIPNs, pseudo-interpenetrating networks, etc. A SIPN is prepared by a process in which the second component

d Unmixed

D

Intercalated

Exfoliated

Figure 1.4 Exfoliation and intercalation of clay nanocomposites.

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5

uniformly swollen

non-swollen

partly swollen

Figure 1.5 Behavior of gels: Uniformly swollen, partly swollen and non-swollen gels.

polymer is polymerized or incorporated following the completion of polymerization of the first component polymer, and thus may be referred to as a sequential SIPN. When an SIPN is prepared by a process in which both component polymers are polymerized concurrently, it may be referred to as a simultaneous SIPN. Pseudo-interpenetrating network systems seem to have interpenetrating networks, but actually do not.

1.3 A Short Survey of the Literature and Applications The 20th century saw great progress in the development and use of polymers and polymer composites; today’s broad family of tailor-made materials allows us to realize the latest technological applications; and the future will see polymers used in increasingly innovative ways. Polymers’ and polymer composites’ almost infinite flexibility and affordability mean that only the imagination of designers limits the ways they can be used. Plastics truly deserve the mantle of material of choice for the 21st century. Just as importantly, polymers’ efficiency allows them to make a real contribution to the vital goals of development for several areas in the world. The growth in these polymeric materials continues to outstrip average annual growth in GDP, and this trend looks set to continue [4]. Demand for polymers in all sectors was up in 2000, continuing the growth trend, and with no significant changes in the relative consumption patterns. Unsurprisingly, the Electrical & Electronic sector shows the highest growth, with countless new inventions and applications using polymers and polymer composites materials as an integral material. However, packaging is still the largest user of plastics, representing about 37% of all other sectors in Europe for 2000. The building and construction industry accounted in 2000 about 19% of total consumption as presented in Figure 1.6 and remains one of the main largest users. The applications in the automotive and electrical & electronics sector are also important and increase year after year.

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Other Household/ Domestic 21.3%

Large Industry 5.4% Agriculture 2.6%

Building and Construction 18.9%

Electrical / Electronic 7.3%

Packaging 37.3%

Figure 1.6 Plastics consumption by industry sector Western Europe 2000. R If we look at the numbers of scientific work published (using the data base ISI Web of Knowledge ) by each country in the last decade with outputs ‘Polymers’ with ‘Blends’, or ‘Gels’, or ‘Nanocomposites’ or ‘Polymer Composites’ we can clearly see in Figure 1.7 that the USA is the leader followed by Asian countries. It is also important to notice that in Western Europe the number of scientific works dedicated to polymers and polymer composites is also important and was the third source of publications in the world.

18000 Multiphase Polymer Systems publications

Publications Intensity 1990-2009

16000 14000 12000 10000 8000 6000 4000 2000 0 Spain

Italy

France England India Germany Japan

China

Others

USA

Country Figure 1.7 Publications intensity in main country on polymers, blends, gels, nano- and macrocomposites between 1990 and 2009.

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Polymers Polymer Composites Polymer Nanocomposites Polymer Blends Polymer Gels C

7

6.98% 2.77% E

D

6.13%

A B

72.6%

11.5%

Figure 1.8 Publications area between 2004 and 2009.

However, it remains that during the last five years (2004–2009) over 72% of the scientific works were dedicated to polymers (Figure 1.8). Polymer composites and polymer nanocomposites materials and blends represent respectively 11.5%, 8% and about 6% of publications, the rest relating to gels.

1.4 Book Content This book is intended to deal with most aspects of multiphase polymer systems science. Four transversal main aspects are considered: 1. The modeling of multiphase polymer systems, including theoretical and numerical simulation approaches. 2. The morphological investigation techniques that allow information on the microstructure of these complex systems to be obtained. 3. The physical characterization at macroscopic scale which also brings information on the structure of the material but mainly determines whether the material properties are compatible with its application and use. 4. The life cycling of multiphase polymer systems covering all the steps between the manufacturing process and the recycling. 1.4.1

Modeling and Computer Simulation of Multiphase Composites: From Nanoscale to Macroscale Properties

Predicting the behavior of multiphase polymer composites is an area with vast potential applications, and has attracted a lot of academic work. Chapters 2, 3 and 4 present the state-of-the-art as well as recent applications. The theoretical approach originates with the Flory-Huggins theory and gives rise to modern powerful predictive methods such as the self-consistent field theory (SCF), which has been applied to the calculation of phase diagrams for diblock copolymer blends, or the study of interfacial properties in binary blends [5, 6]. In Chapter 2 the viscoelastic behavior of polymer nanocomposites in the liquid phase is investigated with recent mesoscale models (sticky reptation model) and the results are quantitatively compared to experiment. The simulation approach is more recent but recent advances make it a very promising field. Modeling can be on the microscale (atomistic), the mesoscale (coarse-grain) or macroscale (continuum)

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Recent work focuses on multi-scale modeling in a ‘bottom-up’ approach where properties obtained at a scale are reused on the larger scale, giving these methods real predictive power. Coarse-grain models and field models are reviewed in detail; dynamical as well as static properties can be investigated with these models. These methods have been applied to miscibility of copolymer blends, or the dynamics of demixing in homopolymer blends. The behavior of polymer multiphase composites is largely influenced by interfacial properties. Chapter 4 proposes a review of interphase properties, interface modification techniques and interface analysis techniques. Mastering the effects of interface properties is a challenge both for modeling and for manufacturing developments [7, 8]. Another important factor in understanding and mastering macroproperties of immiscible polymer composites is morphology (Chapter 6). Two important factors are analyzed in detail: evolution during processing, and component individual properties. Morphology development can be influenced by compatibilization techniques which modify interfacial properties.

1.4.2

Morphological Investigation Techniques

Morphological characterization of the multiphase polymer system is extremely important since most of the physical and transport properties (mechanical, electrical, thermal) of polymer systems are determined by the scale of dispersion of the component phases. The length scale (macro, micro and nano) of the dispersion could be monitored by optical light microscopy (phase contrast, polarized), electron microscopy (SEM, TEM), scanning tunneling microscopy, and atomic force microscopy (AFM). Scattering techniques are also very useful for this; they include light scattering, X-ray scattering (SAXS and WAXS) and neutron scattering techniques. TEM, AFM, SAXS and SANS techniques can also be of interest in the investigation of the interphase/interface of multiphase polymer systems. Several researchers have looked carefully into the interphase/interface width of several multiphase systems using these techniques. Spectroscopy is a valuable technique to characterize multiphase polymer systems. These include UV (ultraviolet) spectroscopy, FTIR (Fourier transform infrared spectroscopy), NMR (nuclear magnetic spectroscopy), XPS (X-ray photoelectron spectroscopy), ESR (electron spin resonance spectroscopy). Each technique varies in its sensitivity. For example, NMR spectroscopy and fluorescence spectroscopy can detect heterogeneities in the range of 2–3 nm scale. NMR can generate a lot of information in the nanoscale on the interphase/interface of multiphase polymer systems. XPS can be used for the study of surface chemistry of polymer grafts, colloidal particles, nanocomposites etc, in the nanometer range.

1.4.3

Macroscopic Physical Characterization

Physical properties of polymers and multicomponents systems based on the use of polymers are more strongly dependent on temperature than for other materials such as metals or ceramics. Moreover, the temperature range of use of these systems is thus reduced when compared to other materials. Physical properties of multiphase polymer systems are also closely related to the components’ properties and relative fractions. The structure of the material – for instance the shape, size, dispersion, possible orientation of the dispersed phase – greatly influences macroscopic properties. For instance, crystallization and melting temperatures of nanocomposites are strongly related to the filler content and its dispersion state whereas the thermal expansion coefficient is strongly affected by the alignment of exfoliated platelets as small changes from perfect planar orientation result in significant changes in thermal expansion behavior.

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9

The properties of interest are of two types: bulk properties and phase-transition properties. The bulk properties are mechanical properties, thermophysical properties, electrical properties or permeability. The two main phase-transition properties concern first-order transitions (fusion and crystallization) and the glass transition. All these properties are now accessible thanks to recent progress in experimental techniques that allow measurements in the three physical states over extended ranges of temperature and pressure, including in the vicinity of the critical point (at least in the case of gases and liquids). The possibility of achieving a complete and reliable characterization of multiphase polymer systems is also one reason for the recent and increasing use of these materials. The knowledge of these properties is important for the development of new systems with properties adapted to one particular application, and also for obtaining information on the microstructure of the material. Thus, macroscopic characterization of multiphase polymer systems is important for engineers in industry and for researchers in the field of material science. Thermophysical properties are important in industry for the knowledge of thermal transport properties and thermal stability of materials. They include thermal conductivity and diffusivity, specific heat, melting and crystallization temperature and enthalpies, coefficient of thermal expansion. Thermal analysis techniques allow information to be obtained on the structure of polymers and multi-components systems and on their phase transitions. Optical techniques allow, for instance, the characterization of the organization of crystalline regions in polymers. Differential Scanning Calorimetry (DSC) is a powerful method for the study of phase transition and for the measurement of specific heat capacity; the recent development of Modulated DSC and the enhanced specifications of recent DSC and MDSC devices have also increased the investigation field of this technique. Chapter 9 will extensively present a survey of thermal analysis applied to the characterization of multi-phase polymer systems. Chapter 10 is devoted to the study of thermophysical properties of multiphase polymer systems (characterization methods, thermophysical behavior, and modeling) [9]. Mechanical behavior is of paramount importance in the design of advanced multiphase polymer materials for many applications in different engineering fields such as aerospace and the automotive industry or civil engineering. The stress–strain behavior, tensile strength, yield strength, elongation at break, hardness, impact behaviors (both notched and unnotched), tear properties, abrasion characteristics and flexural properties are very often determined for the comprehensive understanding of the behavior of multiphase polymer systems. Dynamic mechanical properties are also extremely important for the time-dependent dynamic applications of multiphase polymer systems. Materials reduced to nano-scale can suddenly show very different properties compared to what they exhibit in the macro-scale, enabling unique applications The rheological behavior of multiphase polymer systems is of great importance for the understanding of the flow behavior of these materials. Viscosity-shear rate relationships are fundamental basic data to evaluate the processability of these materials. Although multiphase polymer systems are pseudoplastic, they might show complex behavior such as Newtonian character, yield stress, thixotropic and rheopectic characteristics. Additionally, phase separation, gelation and vitrification can be very well monitored using sensitive rotation rheometers. Very often, rubber modified thermoplastics and other polymer/polymer blend systems show structure build-up at low shear which eventually get destroyed at higher shear forces. The exfoliation/intercalation in polymer nanocomposites could be well understood by careful rheological measurements. Chapters 7 and 8 will be devoted respectively to mechanical characterization and rheology of multiphase polymer systems [10, 11]. The mechanical reinforcement of polymers using nanoparticles will be exposed in Chapter 25. Most polymers behave like electrical insulators. A lot of works was devoted in the last 20 years to the development of highly-conducting polymer systems. Different ways were investigated: synthesis of conducting polymers (such as polypyrrole), mixing of common polymers with metal powders, graphite or, more recently, carbone nanotubes to obtain conducting composites. The challenge is to reach electrical conductivity values minimizing the amount of conducting material mixed with the polymer. The use of conducting nanoparticles as filler in a polymeric matrix is the most efficient solution as very small electrical percolation thresholds are observed in these systems. The study of the electrical behavior of insulating

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polymer systems requires the use of characterization techniques specific to dielectric materials, such as broadband dielectric spectroscopy or thermostimulated depolarization currents. Electrical signals obtained by these techniques are the signature of macromolecular chain movements or of interface effects. The large frequency and temperature ranges covered by these techniques allow a lot of information to be obtained on the microstructure of multiphase polymer systems. Chapter 11 will present a synthesis of literature results concerning the development of conductive polymer systems and of the associated characterization techniques and most commonly used models. Dielectric properties will be investigated in Chapter 12 [12]. Diffusion and transport in multiphase polymer systems have attracted a lot of interest recently (Chapter 19). This is due to the fact that multicomponent polymer systems find enormous application in separation technologies, including dialysis, restricted gas transport, pervaporation, gas sorption and vapor sorption. Several factors such as microstructure, dispersion, miscibility, compatibility, interphase adhesion, degree of crosslinking, orientation of the filler particles, phase co-continuity and so on influence the transport process in multiphase polymer systems. Additionally, diffusion and transport could be used as an excellent tool to characterize the morphology and microstructure of multiphase polymer systems. 1.4.4

Life Cycling

Nowadays, the field of applications of multiphase polymer systems is very broad (civil engineering and buildings, electronics, medicine, transports, packaging. . .). Chapter 23 will give a survey of the existing applications of multiphase polymer systems. Moreover, the continuous development of new materials permits new applications, like fire retardancy applications, which are very important in transport and buildings. This particular topic will be investigated in Chapter 22. Nanoreinforcement of plastics leads to the use of small amounts of fillers to obtain a strong increase of mechanical properties (see Chapters 25 and 6) such as the elastic modulus, even if the use of nanoparticules might cause some health problems. Application allows defining shape and size of the manufactured object and components to be used. All these parameters often impose the use of a particular manufacturing technique. These manufacturing techniques, presented in Chapter 5, are numerous. Classical industrial techniques like extrusion and injection are particularly suitable for the processing of multiphase systems based on the use of thermoplastics. The recent development of new techniques like resin transfer molding has brought new applications concerning processing of fiberreinforced thermoset composites, particularly in the aeronautics and space industries. Finally, the choice of the manufacturing process greatly influences macroscopic properties of the final material (interfaces between components, non-isotropic properties). Therefore, a strong link exists between applications and the manufacturing process. A good knowledge of ageing processes is also required to determine the lifetime of a given object. The study of ageing processes is quite complex as it requires the use of characterization techniques at different length scales and the development of realistic procedures of accelerated ageing (see Chapter 21). Non-destructive testing techniques are of great interest as they can provide an on-line control of the manufactured products and also an early detection of defects, delaminations, etc, during the normal use of a material (see Chapter 20). Recycling of multiphase polymers is a major issue in order to limit the effect of industrial production on the environment. Recycling will be considered in Chapter 24. Possibilities of recycling are mainly dependent on the components used and on the manufacturing process. Thus the possibilities of recycling have to be considered from the beginning of the development of a new material [13].

1.5

Future Outlook, New Challenges and Opportunities

In the area of multiphase polymer systems (blends, composites, nanocomposites, IPNs, gels) several questions still need to be answered. Concerning polymer blends, more sophisticated techniques are needed to probe the

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very early stages of the phase separation process. Although NMR and fluorescence spectroscopy offer some answer to this, we need to have more fast and easy techniques. Similarly for complex systems, we need to look very carefully at the phase separation mechanisms, spinodal or nucleation and growth. There are many complex polymer systems where both mechanisms operate. In the area of compatibilization of polymer blends by reactive route, we need to go a long way to characterize the chemistry of the interfacial chemical reactions. For many polymer–polymer blends systems compatibilized by the reactive route, the chemistry of several interfacial reactions has to be well understood. In order for this we need to go for selective extraction followed by NMR and FTIR spectroscopies. Another important challenge in the area of polymer blends is the online monitoring or in situ monitoring of morphology development during processing. Of course some success has been made in this direction using model systems which are transparent. However, for real polymer systems, we need to undertake more in-depth studies. In this respect, knowledge of the thermophysical properties of polymers over extended ranges of temperature and pressures and in different gaseous environments is undoubtedly necessary to improve the use and lifetime of end-products made of such polymers.In the area of recycling of polymer blends wastes and polymer products, several problems exist which are very important for dealing with waste disposal problems. For effective recycling of polymer blend wastes, we need to have multifunctional compatibilizers which can convert a mixture of a variety of polymer blend materials into useful value-added materials for a number of different applications. However, these techniques have to be more environment-friendly.In the area of composites, the incorporation of naturally-derived macro-, micro- and nanofillers such as cellulose, chitin and starch has recently captured a lot of interest. Since these fillers are derived from waste biomass, the process is highly green and environment-friendly. For polymer nanocomposites, one of the biggest challenges is to reach an excellent dispersion of nanoparticles in the polymer matrix. Although nanoclay is a very efficient reinforcing filler in many polar polymer matrices including nylon, so far we have not been able to achieve good dispersion of the nanoclay in polyolefins (PP, PE) which are one of the plastics with the higher tonnage. More efficient surfactants have to be designed for the excellent dispersion of clay in nonpolar polymer systems such as PP, PE, etc. The extent of intercalation versus exfoliation could not be quantified in many nanoclay-filled polymer systems. We need to have more sophisticated techniques for the exact quantification of the exfoliation process. More environment-friendly and efficient mixing techniques have to be developed for the mixing of nanofillers with various polymer matrices. Of course, strong progress has been made in the use of supercritical carbon dioxide for the processing of various polymer nanocomposites. The orientation of nanoplatelets, such as clay, carbon nanotubes and graphite in the polymer matrix is a major challenge. Perfect alignment of the nanoplatelets will provide excellent properties, particularly gas barrier properties. We need to develop special extrusion techniques, application of magnetic and electrical fields for the orientation of the nanoplatelets. The in situ monitoring of the flow of the polymer nanocomposites during manufacturing requires in-depth research. We still have to go a long way for the successful use of polymer nanocomposites for many commercial applications although some progress has been made in this direction. We should also examine carefully the toxicity aspects of many nanofillers and nanocomposite materials; once the nanofillers enter into our body, elimination will be very difficult. Therefore one has to take extreme care during the handling of the nanofillers in mixing and processing operations. The legal and ethical issues of nanostructured materials have to be addressed carefully. In the area of IPNs, we have made significant progress for the development of nanostructured IPNs. However, the morphology and phase separation mechanisms of nanostructured IPNs have not been well understood. The design of porous networks based on IPNs has received a lot of attention recently. We believe that such porous IPN materials may find potential applications in separation techniques as chromatography supports, as well as membrane catalysis, and more generally in chemistry in confined medium as nanoreactors. However, such new applications have to be explored in detail. There exists much interest for the development of biopolymer-based IPNs for various applications.

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In the case of polymer gels, the development of smart polymer gels with optimum mechanical properties for drug delivery offers enormous opportunites in the medical field. Spatial inhomogeneities present in ionic and nonionic hydrogels depend systematically on network density and on the degree of swelling. The precise quantitative estimation of these inhomogeneities needs more sophisticated and careful analysis. Mechanical instabilities such as buckling, wrinkling, creasing, and folding are commonplace in both natural and synthetic gels over a wide range of length scales. Several advancements have been made on the spontaneous folding behavior of the highly swellable confined nanoscale (thickness below 100 nm) gel films [14]. In fact, the regular self-folding is originated from periodic instabilities (wrinkles) caused by swelling-initiated stresses under confined conditions. Furthermore, folded gel structures can be organized into a regular serpentine-like manner by imposing various boundary conditions on micro-imprinted surfaces. It is important to add that this demonstration of uniform gel to mechanically mediate morphogenesis has far-reaching implications in the creation of complex, large-area, 3D gel nanostructures

References 1. C. W. Macosko, Morphology development and control in immiscible polymer blends, Macromol. Symp. 149, 171–184, 2000. 2. Y. Konishi and M. Cakmak, Structural heirarchy developed in injection molding of nylon 6/clay/carbon black nanocomposites, Polymer, 46, 4811–4826, 2005. 3. N. Nugay and B. Erman, Property optimization in nitrile rubber composites via hybrid filler systems, J. Appl. Polym. Sci., 79, 366–374, 2001. 4. Association of Plastics Manufacturers in Europe, An analysis of plastics consumption and recovery in Western Europe 2000, 2002/GB/04/02, Published Spring 2002. 5. M. Doxastakis, Y. -L. Chen, O. Guzm´an, J. J. de Pablo, Polymer-particle mixtures: depletion and packing effects, J. Chem. Phys., 120, 9335, 2004. 6. M. Matsen and M. Schick, Stable and unstable phases of a diblock copolymer melt, Phys. Rev. Lett., 72, 2660–2663, 1994. 7. G. E. Fantner, O. Rabinovych, G. Schitter, P. Thurner, J. H. Kindt, M. M. Finch, J. C. Weaver, L. S. Golde, D. E. Morse, A. Lipman E, IW, Rangelow, P. K. Hansma, Hierarchical interconnections in the nano-composite material bone: Fibrillar cross-links resist fracture on several length scales, Compos Sci Technol, 66, 1205, 2006. 8. G. Fantner, E. Oroudjev, G. Schitter, L. S. Golde, P. Thurner, M. M. Finch, P. Turner, T. Gutsmann, D. E Morse, H. Hansma, P. K. Hansma, Sacrificial bonds and hidden length: Unraveling molecular mesostructures in tough materials, Biophys J., 90, 1411–418, 2006. 9. R. Tlili, A. Boudenne, V. Cecen, L. Ibos, I. Krupa, Y. Candau, Thermophysical and electrical properties of nanocomposites based on ethylene-vinylacetate copolymer (EVA) filled with expanded and unexpanded graphite, International Journal of Thermophysics, 31(4–5), 936–948, 2010. 10. L. Elias, F. Fenouillot, J. C. Majeste, G. Martin, G., J. Cassagnau, Migration of nanosilica particles in polymer blends, Polym. Sci. : Part B: Polymer Physics 1976–1983, 46, 2008. 11. J. John , D. Klepac, M. Didovi, C. J. Sandesh, Y. Liu, K. V. S. N. Raju, A. Pius, S. Valic and S. Thomas, Main chain and segmental dynamics of semi interpenetrating polymer networks based on polyisoprene and poly(methyl methacrylate), Polymer, 51, 2390–2402, 2010. 12. M. Micusik, M. Omastova, I. Krupa, J. Prokes, P. Pissis, E. Logakis, C. Pandis, P. Potschke and J. Pionteck, A comparative study on the electrical and mechanical behaviour of multi-walled carbon nanotube composites prepared by diluting a masterbatch with various types of polypropylenes, Journal of Applied Polymer Science, 113(4), 2536–2551, 2009. 13. A. Conroy, S. Halliwell and T. Reynolds, Composite recycling in the construction industry, Composites: Part A, 37, 1216–1222, 2006. 14. S. Singamaneni, M. E. McConney and V. V. Tsukruk, Swelling-induced folding in confined nanoscale responsive polymer gels, ACS Nano 4(4), 2327–2337, 2010.

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2 Macro, Micro and Nano Mechanics of Multiphase Polymer Systems Alireza S. Sarvestani Department of Mechanical Engineering, University of Maine, Orono, Maine, USA

Esmaiel Jabbari Department of Chemical Engineering, University of South Carolina, Columbia, USA

2.1 Introduction Flow behavior of colloidal suspensions in complex fluids has been the subject of intense research [1–10]. The widespread use of suspensions in industry has for a long time provided motivation for these investigations. Early theoretical research was based on hydrodynamic models. For example, the low shear rate viscosity of a suspension with low volume fraction, , of solid spherical particles can be estimated by the often-quoted Einstein equation [11] η = ηm (1 + 2.5)

(2.1)

where ηm represents the linear viscosity of the matrix. Underlying these theoretical models is the assumption of continuum viscoelasticity; i.e. the colloidal particle is saturated by the dispersing medium and is large enough such that the non-hydrodynamic contributions such as Brownian motion, surface forces, and van der Waals interactions between particles are negligible. As a result, models derived from hydrodynamic theories are inherently independent from the size of filler particles and the nature of interfacial bonding. It is therefore not surprising that these equations appear to explain only the results of micron-sized particles [5, 12, 13]. They are indeed entirely inadequate to elucidate the results obtained from the reinforcing fillers of colloidal and sub-colloidal size such as silica and graphitized carbon black in non-Newtonian fluids [12, 14, 15].

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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Reduction of the filler size to the order of nanometers can lead to substantial differences in rheology and dynamics of filled polymer liquids compared to complex fluids, reinforced with micron sized particles [14–19]. In particular, polymer composites reinforced with sub-micron fillers, often referred to as polymer nanocomposites, exhibit a significant enhancement in viscoelastic properties compared to microcomposites, at similar filler volume fractions. In frequency sweep tests, this dramatic increase in viscoelasticity is generally manifested at very dilute concentration of filler particles and by the appearance of a secondary plateau for viscoelastic moduli at low frequency regimes [14–16, 20–22]. It has been shown that the mechanical properties of polymer nanocomposites can be greatly modified by changing surface properties of nanoparticles [18, 19, 23]. The extremely large surface area provided by nanoparticles can intensify the effect of particle–particle and/or polymer–particle interactions compared to other reinforcing mechanisms such as hydrodynamic effect. Therefore, the unique properties of polymer nanocomposites are generally rationalized as arising from strong interparticle affinity [24–27] or the interaction between the particle surface and surrounding matrix [28–30]. In the case of strong filler–filler interactions, it is believed that the material response is influenced by the breakdown and reformation of filler agglomerates during mechanical loading [20, 22, 28, 31, 32]. On the other hand, when polymer–filler interactions dominate, it is shown that the suspension viscoelasticity is controlled by the dynamics of stick-slip motion of the polymeric chains on and close to the filler surface [33–37]. Contrary to extensive experimental studies, theoretical models to quantitatively elucidate the reinforcement mechanism in polymer nanocomposites are scarce. Semi-empirical models are proposed for aggregated dispersions based on the concept of fractals [31, 32]. A few rheological models are also presented for conditions where polymer–particle interaction is the dominant reinforcing mechanism [34, 37–39]. The purpose of this chapter is to review our recent progress in constitutive modeling of macroscopic rheological behavior of multiphase polymer systems with strong interaction between fillers and polymer chains (i.e. no filler agglomeration) [40–42]. We have proposed two classes of models corresponding to two different physical regimes: (i) unentangled regime, where the chain length is short (below the entanglement threshold) and topological interactions between chains are unimportant, and (ii) entangled regime, where the chain length is well above the entanglement threshold and the chain motion is severely restricted by topological constraints imposed by entanglement with other macromolecules. Our goal is to identify the relevant physics which control the small-scale structure and kinetic behavior of the composite and apply this information to predict composite macroscopic properties in a series of generic constitutive models. This problem attracts significant practical interest in the field of nanorheology, tribology, heat transfer, and MEMS, due to the promising properties of polymer nanocomposites.

2.2

Unentangled Systems

When the concentration of polymer chains in a suspending medium is below the overlapping concentration, the polymer molecules do not entangle with each other and chain dynamics is governed entirely by friction with solvent or isotropic monomer–monomer friction. For such dilute polymer solutions, generalized beadspring models for the polymer chains can properly describe the kinetics of polymeric liquids [43]. The springs are representative of the elastic entropic tensile forces, while the beads play the role of centers for application of friction forces. The resultant Maxwell-type constitutive equations can provide quantitative description of the concentration and molar mass dependences of terminal relaxation time, terminal modulus and viscosity [43, 44]. In the presence of adhesive filler particles, the influence of reversible filler–polymer interaction (attachment/detachment kinetics) can be easily captured by assuming the attachment point as a region of enhanced friction [38, 39]. This additional friction coefficient is proportional to the corresponding energy of adsorption.

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The structural configuration of adsorbed polymer chains in a good solvent can be described by de Gennes’ theory of reversible adsorption from dilute solutions [45, 46]. It has been shown by many theoretical and experimental studies that the reversible adsorption takes place when the contact energy between the filler surface and each monomer is weak and less than thermal energy, k B T . When the binding energy is somewhat larger than k B T , adsorption becomes irreversible and the chain freezes on the interactive surface [47]. We use the predictions of reversible adsorption theories to study the equilibrium configuration of the polymer layer on the surface of filler particles.

2.2.1

Microscopic Structure

Consider an ensemble of non-entangled flexible polymer molecules and a random distribution of monodisperse non-aggregated rigid spherical particles. The schematic equilibrium configuration of an adsorbed chain on the surface of a particle with radius R f is shown in Figure 2.1. The chain can reversibly adsorb on the colloidal surface and form a polydisperse succession of loops, tails, and sequences of bound monomers (trains). We only consider the case of dilute particle concentration, where no polymer chain may bridge several colloids. Each chain with N monomers (i.e. the lattice sites in Flory-Huggins theory [48]) of size a, occupies a spherical volume with a radius comparable with the Flory radius R F = a N 3/5 in solution. Here, it is assumed that the size of an adsorbed polymer is equal to that of an unperturbed chain in the bulk. This assumption is supported by recent molecular simulation studies for systems with short range (on the scale of one monomer) monomer–surface interactions in the order of k B T [49]. If the interaction range is large and the strength of attraction is relatively high, chains may deform and flatten on the solid surface. This effect is more pronounced for very short chains where the entropic effect is weak [50].

RF

Rf

Figure 2.1 Schematic diagram for equilibrium configuration of an adsorbed polymer chain on a particle surface. The adsorbed chain consists of loops, tails and sequences of bonded monomers.

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The effect of surface curvature on polymer adsorption has been studied by Aubouy and Rapha¨el [51]. Their scaling approach shows that when R f > R F , the curvature of the particle is not relevant and adsorption is similar to that on a flat surface. The adsorption would be enhanced by the effect of surface curvature only at the limit of R F  R f , i.e. when the particles are very small and/or the polymer chains are very long. Although possible, here we do not consider this limiting case, since for solution of non-entangled polymer chains considered in this study, such a small particle size is not practically relevant. In order to describe the structure of the adsorbed and fully equilibrated polymer layer on the filler surface, we use a simplified version of de Gennes scaling theory for reversible adsorption from dilute solutions under good-solvent condition [52]. The configuration of an adsorbed layer is determined by competitive surface attraction and chain entropic interactions. Let us assume that the loops are extended to an average thickness, D, from the filler surface. Since the monomer density is spread over a distance D, we have a f ∼ = D

(2.2)

where f is the fraction of monomers in direct contact with the solid surface. Assuming that the conformational entropy and energetic attraction with the surface are the only factors that determine configuration of the adsorbed layer, the free energy per chain, , can be written as [53] ∼ = kB T



RF D

5/3 − f N E ad

(2.3)

Minimizing the free energy with respect to D yields f ∼ =



E ad kB T

3/2 (2.4)

Equation (2.4) holds true only if the condition of weak monomer–surface coupling is satisfied, i.e. E ad < k B T . Note that even under the limit of weak adsorption energy, the entire chain can be strongly adsorbed due to the many monomer contacts with the surface, i.e. f N E ad  k B T . The chain relaxation dynamics is affected by frictional interactions between monomers and particles. The total friction coefficient as a result of the hydrodynamic force acting on the i th monomer is [39] (ξ )i = ξ1 , (ξ )i = ξ 0 ,

i th monomer is adsorbed th

i monomer is not adsorbed

(2.5a) (2.5b)

where ξ1 is the friction coefficient due to monomer–particle interaction and ξ 0 represents the friction coefficient corresponding to self-diffusion of a single monomer and accounts for its friction with solvent molecules and/or other non-adsorbed monomers. One expects that the surface friction for solid–fluid interfaces to exceed the bulk friction factor in many real systems. Here, it is assumed that this condition is satisfied and hence ξ1 > ξ 0 . Since a fraction f of the monomers in an adsorbed chain is in contact with the particle surface, the total friction coefficient affecting the entire chain is given by ξa = N ( f ξ1 + (1 − f )ξ0 )

(2.6)

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For weakly attractive surfaces, chains are partially adsorbed to the surface and can exhibit their 3-D Rouse dynamics [54]. Considering that the diffusion coefficient of the adsorbed chain is Da = kξBaT , the corresponding relaxation time is ξa τa ∼ = τ f ( f χ + (1 − f )) = R 2F kB T

(2.7)

ξ where τ f ∼ = R 2F k BfT stands for relaxation time of a free chain and χ = ξ10 . According to the self-similar grid structure theory [45], the adsorbed layer can be modeled as a semi-dilute solution of the polymer with continuously varying local concentration of monomers such that at any distance, r , from the surface, the local blob size is equal to r . Therefore, the equilibrium thickness of the layer is on the order of R F . In addition, the equilibrium number of chains in the layer per unit filler surface area is estimated to be ξ

1 0 ∼ = 2 RF

(2.8)

The polymer-filler junctions are transient and their density fluctuates with thermal agitation or the effect of flow. Since the interaction energy is short range, we can assume that the adsorption–desorption process takes place only between those chains which are located at close vicinity of the filler surface with total surface density equal to 0 . If a shows the surface density of adsorbed chains at any instant, the corresponding number of free chains which are able to participate in the adsorption-desorption process is given by p

f =

1 − a R 2F

(2.9)

The rate of attachment and detachment can be shown by the following kinetic equation: ∂ a 1 = ∂t τads



 1 1 − a − a 2 τdes RF

(2.10)

where τads and τdes are the characteristic times of adsorption and desorption of the chains, respectively. The energy expenditure for detachment of an adsorbed polymer molecule is equal to f N E ad . In the presence of an applied deformation rate, the detachment process is favored by the resultant entropic tension exerted by the chains. Considering this effect, the time constants associated with attachment and detachment of the interfacial polymers follow the Arrhenius type relation τdes = A exp τads



f N E ad − δ Fa kB T

 (2.11)

where Fa is the chain entropic force, δ is an activation length on the order of the displacement required to detach the bound chain from the particle surface, and A is a constant. Desorption of a bound monomer with weak and short-range interaction with the adsorbing surface can be considered as a local process. This takes place by diffusing a distance on the order of the equilibrium size of the first blob in contact with the wall [55]. According to the self-similar grid structure theory [45], the size of the first blob in contact with the particle surface is on the order of the monomer size. Therefore, δ, the total displacement required to separate the entire chain with f fraction of adsorbed monomers is a ≤ δ ≤ R F .

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Macroscopic Properties

The classical Maxwell model is used to describe viscoelasticity of the composite [44]. At any instant, a representative chain is either adsorbed to the particle surface or it is free. The total volume average stress in the polymer matrix is therefore σ = ψσ a + (1 − ψ)σ f

(2.12)

where σ a and σ f are the stress contribution of the adsorbed and free chains, respectively, and ψ shows the volume fraction of the polymer–particle interfacial zone (with thickness ∼ R F ) in the matrix. Since the particles are homogeneously dispersed, all statistical properties corresponding to any arbitrary representative mesodomain of the composite body are assumed to be statistically homogeneous. Hence, the multipoint statistical moments of any order are shift-invariant functions of spatial variables for any ergodic field and, therefore, the ensemble averaging could be replaced by volume averaging [56]. The contribution of polymer chains to the stress tensor is given by Kramers expression σ a = 3k B T Na σ f = 3k B T N f

Ra Ra  R2  F  Rf Rf

(2.13a) (2.13b)

R 2F

where Na and N f represent the number density of the adsorbed and free chains, respectively. Ri (i = a, f ) is the chain end-to-end vector and · · · shows the ensemble average. Assuming that the number density and surface density of interfacial chains follow the relation N ∼ = /R F , in steady state situation, we can write Na p = A exp Nf



f N E ad − δ Fa kB T

 (2.14)

p

where N f shows the number density of free chains within the interphase zone. In their simplest forms, the constitutive relations for the evolution of rate dependent stresses produced by the chains can be expressed by Maxwell (upper-convected) equations τa σˆ av + σ av − G a I = 0 τ f σˆ vf + σ vf − G f I = 0

(2.15a) (2.15b)

where I is the identity tensor and G i is the stiffness (i = a, f ). Here, σˆ designates the upper-convected − σ · L e f − L Te f .σ , where L e f = h()∇v is the effective derivative of the stress tensor given by σˆ = ∂σ ∂t velocity gradient tensor and v is the velocity field. Here, h() accounts for the hydrodynamic interaction between particles with volume fraction . The contribution of the hydrodynamic effect is controlled by shape and volume fraction of the particles [11]. At low filler concentrations, it is typically represented by h() = 1 + ζ  where pre-factor ζ accounts for the geometry of particles.

(2.16)

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2.2.3

19

Results

In this section, the proposed model is used to predict the steady state response of a polymer solution under 2-D shear flow, reinforced with spherical rigid particles (ζ = 2.5). The velocity gradient tensor in this case can be represented by ⎛

0 ∇v = ⎝ 0 0

⎞ 0 0⎠ 0

γ˙ 0 0

(2.17)

where γ˙ is the shear rate. In our parametric study, R F (Flory radius) and τ f (relaxation time of free chains) are taken to be the unit of length and time, respectively. The activation length δ is assumed to be on the order of ∼0.5R F and the dimensionless constant A is set equal to 10-4 . We assume that each chain is comprised from 5000 monomers and the concentration of dilute polymer solution is ∼0.1(R −3 F ). The energetic affinity between the polymer R ad and ρ = R Ff , chains and fillers and size of the fillers are characterized by dimensionless quantities ε = E kB T respectively. Figure 2.2(a) represents the dependence of low shear rate steady state (LSRSS) viscosity of the reinforced polymer on the polymer–particle interaction through affinity parameter ε. The values are normalized to the steady state viscosity of the unfilled solutions (ηm ). Here, it is assumed that  = 10% and χ = 10. The results clearly show the strong effect of monomer-filler energetic affinity on the overall shear viscosity of the mixture, especially within the range of small particle size. Figure 2.2(b) shows the size effect of spherical fillers on LSRSS viscosity of the composites at  = 10% and χ = 10. Even within the range of weak monomer-particle attraction considered in this chapter, the concentration of adsorbed chains at the vicinity of particles could be substantially higher than that in the bulk. This indicates an indirect dependence of the

(b)

(a) 60

Normalized LSRSS viscosity

ρ=5 ρ=1

50

100

Φ = 0.1 χ = 10

Normalized LSRSS viscosity

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0.1

0.2

0.3

0.4

ε = Δ Ead / kBT

0.5

0.6

0.7

Φ = 0.1 χ = 10

ε = 0.1 ε = 0.5

10

1 1

10

100

ρ = R f /R F

Figure 2.2 The variation of LSRSS viscosity with (a) monomer–particle interaction energy (ε = E ad /kB T ) and (b) the filler size (ρ = R f /R F ). Volume fraction of dispersed particles was assumed to be 10% ( = 0.1) and relative monomer–particle friction is taken to be χ = ξ1 /ξ0 = 10 in all cases. The viscosity values are normalized to the viscosity of unfilled polymer solution.

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10 Normalized LSRSS viscosity

9 8 7 6

(b) 9

Φ = 0.01 Φ = 0.05 Φ = 0.10

χ = 10 ρ = 5 ε = 0 3

Φ = 0.15 Φ = 0.20

5 4 3 2 1 0 10

ρ = 50

8 Normalized LSRSS viscosity

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ρ=5

7 6 5 4 3 2 1 0

γτf

100

0

0.05

0.1

0.15

0.2

0.25

Filler volume fraction (Φ)

Figure 2.3 (a) The variation of LSRSS viscosity with shear rate at different filler concentrations (). (b) Effect of filler size (ρ = R f /R F ) on the variation of viscosity with filler concentration.

overall properties on filler size. At the limit of small particles, this high density interfacial layer achieves a relatively high volume fraction and can significantly affect the steady state shear viscosity of the mixture. By increasing the particle size, the overall viscosity of the composite converges to h()ηm , implying that the hydrodynamic interaction is the dominant reinforcing mechanism in composites made with larger particles. The variation of steady state viscosity of the suspensions with applied shear rate is depicted in Figure 2.3(a) for different concentration of filler particles. All other parameters are kept constant. The neat dilute polymer solution is considered to be shear rate independent and, hence, the observed nonlinearity is totally due to the thixotropic effect of chain-filler detachment. The detachment process is favored by the resultant entropic tension exerted by the chains which, according to Eq. (2.14), leads to a reduction of the number density of adsorbed chains at higher deformation rates, and a significant decrease in shear viscosity of the filled systems. Figure 2.3(b) shows the values of LSRSS viscosities as a function of filler volume fraction, , at two different filler sizes. In the case of ρ = 5, the shear viscosities are almost equal to the predictions of Eq. (2.1). However, when the particle size is comparable with molecular dimensions (i.e. ρ = 1) the magnitudes of shear viscosity are significantly higher than the predictions of Eq. (2.1). This phenomenon is experimentally observed for suspensions of well-dispersed nanoparticles at low concentrations [14]. In the present model, this behavior is solely due to the energetic interaction between polymer chains and fillers. Figure 2.4(a) demonstrates the effect of χ on the storage modulus G  (ω) in a frequency sweep, where the values of , ρ, and ε are fixed. When the monomer-surface friction is weak, the typical slope of the neat polymer (2:1 in log-log scale) is obtained for the filled system at low frequencies. As the ratio ξ1 /ξ0 increases, a secondary plateau in the low frequency region forms which is associated with the relaxation of adsorbed chains. This transition, from liquid to solid-like behavior, indicates that the stress relaxation can be effectively hindered by the presence of nanoparticles, when the relaxation time and number density of the adsorbed chains are sufficiently high. This deceleration in relaxation process at low frequencies is reported frequently for a variety of nanofilled polymer suspensions. Figure 2.4(b,c) shows a similar behavior corresponding with the increase in polymer–particle energetic affinity and reduction of particle size, both of which result in enhancement of the concentration of adsorbed chains in the system. The results of frequency sweeps, for suspensions at different filler concentrations, are also shown in Figure 2.4(d), where all other parameters are

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(a) 1

1

χ = 10

ρ=5

χ = 50 0.1

G′ k BT

ρ = 20

χ = 1000

0.1

G′ k BT

0.01

ρ = 50

0.01

0.001

0.001

Φ = 0.01 χ = 100

Φ = 0.01 ε = 0.5 0.0001 0.01

0.1

ωτf

1

0.0001 0.01

10

0.1

1

ωτf

10

(d)

(b) 1

0.1

21

1

ε = 0.1

Φ = 0.01

ε = 0.3

Φ = 0.05

ε = 0.5

Φ = 0.10 0.1

G′ k BT

G′ k BT

0.01

0.01

0.001

ε = 0.5

Φ = 0.01

χ = 100

χ = 100 0.0001 0.01

0.1

ωτf

1

10

0.001 0.01

0.1

ωτf

1

10

Figure 2.4 The variation of frequency response of the filled suspensions with (a) friction between the monomers and particle (χ = ξ1 /ξ0 ), (b) monomer–particle interaction energy (ε = E ad /kB T ), (c) filler size (ρ = R f /R F ), and (d) filler volume fraction ().

kept constant. This figure shows that at higher concentration of dispersed particles, the system could show a relatively independent frequency behavior. This rubber-like behavior, which is also reported experimentally [15, 20, 21], is due to the increase in volume fraction of the polymer–filler interfacial zone (especially for smaller filler particles) and hence the increase in fraction of adsorbed chains in filled systems.

2.3 Entangled Systems In polymer melt of sufficiently long chains or in polymer solution of sufficiently high concentration (above the overlap concentration), the dynamics of flexible polymer chains is controlled by the effect of topological

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constraints imposed by interchain entanglements [43]. Since the covalently bonded molecular chains cannot cross each other, the large-scale mobility of an entangled chain is restricted to its one-dimensional Brownian motion along the backbone. Hence, the overall viscoelastic behavior of entangled systems appears to depart significantly from the predictions of general bead-spring models. Instead, the well-known reptation model of de Gennes, Doi and Edwards [43] can more realistically represent the influence of topological constraints on diffusional motion of entangled polymers. Obviously, adhesive particles dispersed in an entangled suspending medium reduce chain mobility. However, due to the reversible nature of energetic interaction between fillers and chains, the reptative motion of the chains is not fully suppressed; adsorbed chains have the possibility to exercise their reptative motion within the encompassing tube during those time intervals which are simultaneously detached from adsorbing particles. Full relaxation is reached when the representative chain completely renews its original tube. The effective reptation time of the reversibly adsorbing chain can be determined using the sticky reptation model [42].

2.3.1

Microscopic Structure

Consider an ensemble of entangled polymer molecules with Ne entanglement segments and a random distribution of non-aggregated rigid spherical nanoparticles. The monodispersity of both species is assumed. The schematic configuration of an adsorbed chain is shown in Figure 2.1. The chains can reversibly adsorb on the colloidal surface and form a polydisperse succession of loops, tails and sequences of bound monomers. It is also assumed that the dispersed particles are sufficiently small such that even at low volume fractions, the average particle wall-to-wall distance could be on the order of the average size of a polymer coil. Hence, at equilibrium, chains may simultaneously attach to more than one nanoparticle. The effect of topological constraints on an adsorbed chain can be modeled by a confining tube with average diameter a, equal to the size of an entanglement segment (Figure 2.1). It has been suggested that the effective entanglement length may change close to an impenetrable wall [57]. In that case, this is a local effect which decays on a length scale comparable to the bulk value of a. In the present analysis, this possible effect is neglected and the tube diameter of adsorbed chains is taken to be similar to that of free chains. The probability for an entangled chain to make contact with nanoparticles at Nc entanglement segments can be approximated by  p(Nc ) =

Ne Nc



q Nc (1 − q) Ne −N c

(2.18)

where q is the probability for one entanglement segment to bind to the surface of a nanoparticle. The underlying assumption for Eq. (2.18) is that the number of adsorbed entanglement segments of a single macromolecule is a random quantity and all entanglement segments have equal probability to attach to a nanoparticle. Obviously, the accuracy of this approximation for the probability distribution of Nc depends on the spatial distribution of nanoparticles and the relative values of Nc and Ne . Given the fact that q is usually a very small number, the distribution function shown by Eq. (2.18) can be approximated by

p(Nc ) =

Ne ! q Nc exp[−q(Ne − Nc )] Nc !(Ne − Nc )!

(2.19)

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A B

C a

D Figure 2.5 Schematic configuration of a free and an adsorbed chain in an entangled melt. The effect of molecular entanglement of surrounding chains is represented by encompassing tubes with diameter a. The adsorbed chain is attached on the surface of two nanoparticles at points B and C.

For a long polymer chain, well above the entanglement threshold, Ne is usually very large. At the limit of Ne  Nc , we have 1 (N c ) Nc exp[−N c ] p(Nc ) ∼ = Nc !

(2.20)

where N c = Ne q is the average number of adsorbed entanglement segments per chain. In addition to the adsorbed chains, there is a population of free chains, which are not interacting with any particle (during the characteristic time of their relaxation), and hence simply exercise their regular reptative motion (Figure 2.5). The free reptation of the adsorbed chains, however, is suppressed because of their association with nanoparticles. Since the polymer–particle interaction is considered to be a reversible process, the adsorbed entanglement segment desorbs from the particle surface after a finite residence time τ , due to thermal fluctuation. This quantity can be utilized to express another definition for N c in Eq. (2.3) as N c = ρτ , where ρ represents the average number of attached entanglement segments in a chain per unit time. If each nanoparticle is considered as a temporary crosslink point between adsorbing entangled polymers, then the diffusion of adsorbed chains can be considered to be somewhat similar to the diffusion of associating polymers [58, 59]. Within this framework, the relaxation of the entire adsorbed chain is the result of partial relaxation of bridged segments (e.g. BC in Figure 2.5) and dangling ends (e.g. AB or AC after detachment of the chain from the particle at B) due to longitudinal Rouse type motion of these segments along the encompassing tube, in addition to the reptative motion of the entire chain during the time intervals in which the chain is not attached to any nanoparticle. Here, as a first order approximation, we neglect the partial relaxation of bridged segments and dangling ends and assume that an adsorbed chain can only make a reptation step when it is simultaneously detached from all temporary junctions. This relaxation process is called sticky reptation. Implementation of the relaxation of bridges and dangling ends in the model is possible; however, this is not examined here in order to reduce the number of adjustable parameters when the model predictions are compared with experiments.

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Let us assume that tad is the average time in which a chain stays in contact with nanoparticle(s). Each tad is followed by a time interval equal to tfr that chain is free and carries no temporal crosslink. According to Eq. (2.20), the fraction of time that a chain spends free can be evaluated as tfr = p(0) ∼ = exp[−Ne q] tad + tfr

(2.21)

We are interested in estimating the average reptation time for adsorbed chains. This represents the characteristic time that the chain needs to completely renew its encompassing tube. Since it is assumed that the adsorbed chains can only diffuse during tfr , then their effective reptation time, τa , can be estimated by [42] τf ∼ τa ∼ = = τ f exp[Ne q] p(0)

(2.22)

where τ f represents the characteristic reptation time of free chains. Parameter q implicitly depends on the volume fraction of dispersed nanoparticles, their size, and their interaction with surrounding polymer molecules.

2.3.2

Macroscopic Properties

Equation (2.22) shows that by increasing q, the relaxation of adsorbed chains could be considerably slower than the free chains. Therefore, for a system composed of populations of macromolecules with fast and slow relaxation dynamics, the mechanism of stress relaxation seems to be similar to double reptation of binary blends, due to the effect of constraint release [33–36]. This possibility is examined in the next section of the chapter. Here, we express the stress relaxation function in the following form to include chain reputation as well as tube renewal due to constraint release of the surrounding chains [60]: G(t) = G 0 [(1 − ψ)μ f (t) + ψμa (t)]α

(2.23)

where φ is volume fraction of the adsorbed chains and G 0 is plateau modulus. μ f and μa represent the fraction of surviving tube segments for free and adsorbed chains, respectively, which are defined by [61]:   8 1 −tn 2 exp π 2 n n2 τf

(2.24a)

  8 1 −tn 2 μa (t) = 2 exp π n n2 τa

(2.24b)

μ f (t) =

odd

odd

Parameter α in Eq. (2.6) represents the contribution of constraint release in stress relaxation. If it is assumed that each chain moves inside its tube, independent of the motion of other chains (i.e. no constraint release), then α = 1, following the Doi-Edwards version of the reptation model. If the diffusional motion of surrounding chains is taken into account, then it can be shown that α = 2, following the theory of double

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25

(b)

1

1

0.1

0.1

G′/G 0

G ′/G 0 ψ=0

0.01

ψ = 0.1

0.01

c = 10 c = 50

ψ = 0.2

c = 100

ψ = 0.5 0.001 0.01

0.1

1

10

c = 500 0.001 0.01

ωτ f

0.1

1

10

ωτ f

Figure 2.6 The variation of normalized storage modulus with (a) concentration of adsorbed chains (ψ) and (b) relative relaxation times of adsorbed and free chains (c = τa /τ f ) in a frequency sweep.

reptation [60, 62, 63]. The storage and loss moduli as functions of frequency ω can be calculated from Eq. (2.23), using the following expressions: 

∞

G (ω) = ω

G(t) sin(ωt)dt

(2.25a)

G(t) cos(ωt)dt

(2.25b)

0

G  (ω) = ω

∞ 0

2.3.3

Results

We start with a parametric study on the effect of fraction and relaxation time of the adsorbed chains on the linear (steady state) viscoelastic response of a filled melt. Here, only the principal mode of relaxation is taken into account (i.e. n = 1 in Eq. (2.24)). The relative relaxation time of the adsorbed and free chains is characterized by c = τa /τ f and the parameter α is set equal to unity (i.e. no constraint release). Figure 2.6(a) represents the effect of concentration of adsorbed chains (ψ) on the normalized shear modulus G  in a frequency sweep. The parameter c is taken to be constant and equal to 100. Unfilled systems (ψ = 0) show the typical slope of 2:1 (in log-log scale) at low frequencies, providing a comparison for the effect of chain adsorption. As the concentration of adsorbed chains increases, the magnitude of storage modulus increases and a secondary plateau in low frequency region forms. The emergence of the plateau region is manifested when only 10% of the polymer chains are assumed to be adsorbed. Figure 2.6(b) shows that increasing the relaxation time of the adsorbed chains (at constant concentration of 20%) leads to a similar trend in viscoelastic behavior. These results indicate that the model is able to predict the enhancement in elasticity and formation of secondary plateau at low frequency domains, similar to the frequency sweep results of nanofilled polymer melts. According to the model, this behavior is associated with the increase in fraction and relaxation time

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of the adsorbed chains. The magnitude of these quantities is presumably controlled by the volume fraction and size of the dispersed particles, as well as the interaction between particles and surrounding chains. The surface area of well-dispersed colloidal particles increases by raising their volume fraction and more effectively by reducing their diameter. These effects increase the probability of polymer adsorption (i.e. q) and the fraction of adsorbed chains (i.e. ψ). Moreover, in the presence of strong polymer–particle interactions, the average residence time τ of the entanglement segments near the particle surface increases, which results in a significant reduction in mobility of the adsorbed chain. In the presented theory of sticky reptation, this effect is shown by the exponential dependence of relaxation time τa with quantity N c = ρτ , as indicated by Eq. (2.22). This implies that the stress relaxation at low frequencies (comparable with τa−1 ) can be hindered by the presence of nanoparticles, leading to a transition from liquid to solid-like behavior in those frequency regions. This deceleration in relaxation process at low frequencies is reported frequently for a variety of nanofilled polymer melts. Zhong and Archer [21] have reported on the linear viscoelastic response of poly(ethylene oxide) (PEO) with nearly monodisperse molecular weight of 189 kDa (i.e. N = 86Ne ), reinforced with isotropic silica nanospheres with average diameter of 12 nm. They have conducted relaxation and frequency sweep tests on PEO/silica nanocomposites in the melt state with different concentrations of dispersed nanoparticles. In what follows, the predictions of the proposed model in this chapter are quantitatively compared with their experimental data. The fitting parameters of the study include G 0 , τ f , q, and ψ. The values of the first three parameters extracted from fitting with experiments are 1300 kPa, 0.7 sec, and 0.084. ψ is the only fitting parameter whose value changes with the concentration of silica nanoparticles; we found that ψ = 0.15 at 2 vol% silica whereas ψ = 0.45 at 4 vol% of silica. Figure 2.7 compares the model predictions and experimental data for G  in frequency sweeps. The best fits to the data of unfilled melts were obtained when the effect of constraint release was taken into account (i.e. α = 2). In contrast, predictions are in much better agreement

1.0E+07

1.0E+06

1.0E+05

1.0E+04

G ′ (Pa)

no silica 2 vol% silica 4 vol% silica

1.0E+03

1.0E+02

1.0E+01

0.01

0.1

1

10

100

ω(1/s) Figure 2.7 Comparison of the predicted storage modulus by model (solid lines) with experimental results of Zhong and Archer. Reprinted from [21]. (12/24/2002) Copyright (2002) American Chemical Society.

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27

with experimental data of filled melts when α = 1 and the effect of constraint release is neglected. This can be attributed to the relatively frozen dynamics of adsorbed chains in filled systems. In this way, a fraction of the tube forming chains exhibits an extremely slow relaxation and, as a result, the survival time of the encompassing tube increases. By increasing the concentration of nanoparticles and consequently the fraction of adsorbed chains, an entrapped chain exercises its reptation motion in a practically time-independent network and without constraint release. Further investigation, however, is necessary to validate the generality of this observation for different nanocomposites melts.

2.4 Conclusion Growing applications of polymer nanocomposites require a comprehensive understanding of their equilibrium and kinetic structure at a wide spectrum of time and length scales. Different techniques at various time and length scales from molecular scale (e.g. atoms), to mesoscale (e.g. coarse-grains, particles, monomers) and to macroscale, have shown success in addressing various aspects of polymer nanocomposite properties. In this chapter, we reviewed two classes of mesoscale models for linear viscolesatic behavior of unentangled polymer solutions and entangled polymer melts reinforced with non-agglomerated nanoparticles at low volume fraction and with short range energetic affinity with polymer chains. In filled dilute polymer solutions, the structure of the adsorbed polymer layer was determined using the scaling theory of adsorption in good-solvent conditions. The dynamics of the system was modeled by the classical Maxwell constitutive relations. The relaxation of entangled chains was described by the sticky reptation model [42]. The system relaxation was analyzed by the combination of stress relaxation of free and adsorbed chains. The proposed models were able to rationalize the specific properties of reinforced polymer liquids based on the role of energetic affinity towards particles at sub-colloidal sizes. This effect intensifies by increasing volume fraction and reducing size of the particles which leads to higher fraction of adsorbed chains in the system. It was shown that behaviors such as high shear viscosity at low filler concentrations or solid-like properties in low frequency regimes could be attributed to the slowdown of the relaxation process in polymer chains. This process is controlled by the frictional interaction between monomers and particles, the density of the adsorbed polymer chains on the particle surface (controlled by monomer-particle adsorption energy), and volume fraction of the interfacial layer which can be enhanced by reduction of filler size or increasing filler concentration. Although the proposed models show the ability to predict different characteristics of viscoelastic properties of nanofilled polymers, they have a few notable limitations. The models hold true only for systems with homogeneously dispersed nanoparticles. This is an idealization, since the formation of particle aggregates is always possible in experimental conditions (due to either incomplete exfoliation during synthesis or deformation-induced flocculation). The porosity in the structure of individual aggregates increases the effective volume fraction of the rigid phase, which can lead to a significant contribution of hydrodynamic effect. In addition, the clusters of aggregated particles follow a different relaxation pattern which is rooted in the stored elastic energy of the strained clusters and the failure/restoration properties of filler–filler bonds. Moreover, in the proposed model for entangled systems, the effects of filler size and concentration are only implicitly accounted for. Understanding how the values of q and φ change with size and concentration of nanoparticles is possible by performing complementary molecular simulation studies on relevant systems. Finally, it should be emphasized that in the proposed models, the adsorption of molecules on the surface of nanoparticles is regarded as a reversible phenomenon. In reality, this assumption is not necessarily valid for any filled system. For a polymer melt, a subtle interplay between compressibility, monomer energetic affinity with the surface, and surface/chemical heterogeneities is expected to determine reversibility or irreversibility of polymer–particle interaction [64].

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Acknowledgements A.S. Sarvestani gratefully acknowledges the financial support provided by the Department of Mechanical Engineering and Office of Vice President for Research at the University of Maine. This work was supported by a grant to E. Jabbari from the National Science Foundation under Grant No. CBET-0756394.

References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39.

I. M. Krieger, Adv. Colloid Interface Sci. 1972, 3, 111. D. J. Jeffrey, A. Acrivos, AIChE J. 1976, 22, 417. J. Mewis, A. J. B. Spaull, Adv. Colloid Interface Sci. 1976, 6, 173. C. E. Chaffey, Colloid Polym. Sci. 1977, 255, 691. A.B. Metzner, J. Rheol. 1985, 29, 739. W. B. Russel, D. A. Saville, W. R. Schowalter, “Colloidal Dispersions”, Cambridge University, London 1989. P. J. Carreau, P. A. Lavoie, M. Bagassi, Macromol. Symp. 1996, 108, 111. H. A. Barnes, J. F. Hutton, K. Walters, “An Introduction to Rheology” Elsevier, Amsterdam 1989. H. A. Barnes, J. Non Newton. Fluid Mech. 1997, 70, 1. K. J. Kim, J. L. White, Polym. Eng. Sci. 1999, 39, 2189. A. Einstein, Ann. Phys. 1906, 19, 289. G. Kraus, Rubber Chem. Teehnol. 1965, 38, 1070. A. J. Poslinski, M. E. Ryan, R. K. Gupta, S. G. Seshadri, F. J. Frechette, J. Rheol. 1988, 32, 704. Q. Zhang, L. A. Archer, Macromolecules 2004, 37, 1928. T. Zhu, T. Thompson, S. Q. Wang, E. D. von Meerwall and A. Halasa, Macromolecules 2005, 38, 8816. J. F. Le Meins, P. Moldenaers, J. Mewis, Ind. Eng. Chem. Res. 2002, 41, 6297. O. Harzallah, D. Dupuis, Rheol. Acta 2003, 42, 10. M. A. Osman, A. Atallah, Polymer 2005, 46, 9476. M. A. Osman, A. Atallah, Polymer 2006, 47, 2357. M. I. Aranguren, E. M. Mora, J. V. DeGroot, C. W. Macosko, J. Rheol. 1992, 36, 1165. Q. Zhang, L. A. Archer, Langmuir 2002, 18, 10435. P. Cassagnau, F. M´elis, Polymer 2003, 44, 6607. ¨ M. A. Osman, A. Atallah, T. Schweizer, H. C. Ottinger, J. Rheol. 2004, 48, 1167. A. Medalia, J. Colloid. Interface Sci. 1970, 32, 115. J. Fr¨olich, S. Kreitmeier, D. G¨oritz, Kautsch. Gummi Kunstst. 1998, 51, 370. J. M. Piau, M. Dorget, J. F. Palierne, A. Pouchelon, J. Rheol. 1999, 43, 305. D. J. Kohl, G. Beaucage, Curr. Opin. Solid State Mater Sci. 2002, 6, 183. J. Rehner, in: Reinforcement of Elastomers, G. Kraus, Ed., Interscience, New York 1965, p. 153. J. Leblanc, J. Appl. Polymer Sci. 1997, 66, 2257. M. Wang, Rubber Chem. Tech. 1998, 71, 521. T. A. Witten, M. Rubinstein, and R. H. Colby, J. Phys II 1993, 3, 367. W. Wolthers, D. van den Ende, V. Breedveld, M. H. G. Duits, A. Potanin, R. H. W. Wientjes, J. Mellema, Phys. Rev. E 1999, 56, 5726. M. Simhambhatla, A. I. Leonov. Rheol. Acta 1995, 34, 329. G. Havet, A. I. Isayev, Rheol Acta 2001, 40, 570. G. Havet, A. I. Isayev, Rheol. Acta 2003, 42, 47. H. Montes, F. Lequeux, J. Berriot, Macromolecules 2003, 36, 8107. V. Pryamitsyn, V. Ganesan, J. Rheol. 2006, 50, 655. A. S. Sarvestani, C. R. Picu, Polymer 2004, 45, 7779. A. S. Sarvestani, C. R. Picu, Rheol. Acta 2005, 45, 132.

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Macro, Micro and Nano Mechanics of Multiphase Polymer Systems 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57. 58. 59. 60. 61. 62. 63. 64.

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A.S. Sarvestani, E. Jabbari, Biomacromolecules 2006, 7, 1573. A. S. Sarvestani, E. Jabbari, Macromol. Theory Simul. 2007, 16, 378. A. S. Sarvestani, Eur. Polym. J. 2008, 44, 263. M. Rubinstein, R.H. Colby, “Polymer Physics”, Oxford University Press, London 2003. G. Larson, “ Constitutive Equations for Polymer Melts and Solutions”, Butterworths, Boston 1988. P. G. de Gennes, Macromolecules 1981, 14, 1637. P. G. de Gennes, Adv. Colloid Interface Sci. 1987, 27, 189. B. O’Shaughnessy, D. Vavylonis, J. Phys. Condens. Matter. 2005, 17, R63. P. J. Flory, “Principles of Polymer Chemistry”, Cornell University Press, New York 1953. M. S. Ozmusul, C. R. Picu, Polymer 2002, 43, 4657. M. Doxastakis, Y. -L. Chen, O. Guzm´an, J. J. de Pablo, J. Chem. Phys. 2004, 120, 9335. M. Aubouy, E. Rapha¨el, Macromolecules 1998, 31, 4357. P. G. de Gennes, “Scaling Concepts in Polymer Physics”, Cornell University Press, Ithaca 1985. J. Gong, Y. Osada, J. Chem. Phys. 1998, 109, 8062. A. L. Ponomarev, T. D. Sewell, C. J. Durning, J. Polym. Sci. Polym. Phys. 2000, 38, 1146. J. Wittmer, A. Johner, J. -F. Joanny, K. Binder, J. Chem. Phys. 1994, 101, 4379. E. Kr¨oner, “Statistical Continuum Mechanics”, Springer Verlag, Wien 1972. X. Zheng, B. B. Sauer, J.G. van Alsten, S. A. Schwarz, M. H. Rafailovich, J. Sokolov, M. Rubinstein, Phys. Rev. Lett. 1995, 74, 407. A. E. Gonzalez, Polymer 1983, 24, 77. L. Leibler, M. Rubinstein, R.H. Colby, Macromolecules 1991, 24, 4701. M. Marrucci, J. Polym. Sci. Polym. Phys. Ed. 1985, 23, 159. M. Doi, S.F. Edwards, “The theory of polymer dynamics”, Oxford, Clarendon 1986. J. des Cloizeaux, Macromolecules 1990, 23, 3992. J. des Cloizeaux, Macromolecules 1990, 23, 4678. G. D. Smith, D, Bedrov, O, Borodin, Phys. Rev. Lett. 2003, 90, 226103.

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3 Theory and Simulation of Multiphase Polymer Systems Friederike Schmid Institute of Physics, Johannes-Gutenberg Universit¨at Mainz, Germany

3.1 Introduction The theory of multiphase polymer systems has a venerable tradition. The ‘classical’ theory of polymer demixing, the Flory-Huggins theory, was developed in the 1940s [1, 2]. It is still the starting point for most current approaches – be they improved theories for polymer (im)miscibility that take into account the microscopic structure of blends more accurately, or sophisticated field theories that allow to study inhomogeneous multicomponent systems of polymers with arbitrary architectures in arbitrary geometries. In contrast, simulations of multiphase polymer systems are quite young. They are still limited by the fact that one must simulate a large number of large molecules in order to obtain meaningful results. Both powerful computers and smart modeling and simulation approaches are necessary to overcome this problem. In the limited space of this chapter, we can only give a taste of the state-of-the-art in both areas, theory and simulation. Since the theory has reached a fairly mature stage by now, many aspects of it are covered in textbooks on polymer physics [3–9]. The information on the state-of-the art of simulations is much more scattered. This is why we have put some effort into putting together a representative list of references in this area – which is of course still far from complete. The chapter is organized as follows. In Section II, we briefly introduce some basic concepts of polymer theory. The purpose of this part is to make the chapter accessible to readers who are not very familiar with polymer physics; it can safely be skipped by the others. Section III is devoted to the theory of multiphase polymer systems. We recapitulate the Flory-Huggins theory and introduce in particular the concept of the Flory interaction parameter (the χ parameter), which is a central ingredient in most theoretical descriptions of multicomponent polymer systems. Then we focus on one of the most successful mean-field theories for inhomogeneous (co)polymer blends, the self-consistent field theory. We sketch the main idea, discuss various

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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aspects of the theory and finally derive popular analytical approximations for weakly and strongly segregated blends (the random phase approximation and the strong-segregation theory). In Section IV, we turn to discussing simulations of multiphase polymer systems. A central concept in this research area is ‘multiscale modeling’: Polymers cannot be treated at all levels of detail simultaneously, hence coarse-grained models are used in order to study different aspects of the materials in different simulations. This allows one to push the simulation limits to larger length and time scales. We describe some of the most popular coarse-grained structural and dynamical models and give an overview over the state-of-the-art of simulations of polymer blends and copolymer melts.

3.2

Basic Concepts of Polymer Theory

For the sake of readers who are not familiar with polymer theory, we begin with recapitulating very briefly some basic concepts. Polymers are macromolecules containing up to hundreds of thousands of atoms. At first sight, one would not expect such molecules to be easily amenable to theoretical modeling; however, it turns out that the large size of the molecules and their highly repetitive structure in fact simplifies things considerably. Since polymer molecules interact with many others, details of local interactions average out and polymers can often be characterized by a few effective quantities, such as their topology, the local stiffness along the backbone, the bulkiness, the compatibility/incompatibility of the building blocks, etc. Decades ago, pioneers like Flory [3], Edwards [5] and de Gennes [4] have established theoretical polymer science as a highly successful field of research, which brings together scientists from theoretical chemistry, statistical physics, materials science, and even the biosciences, has created a wealth of new beautiful theoretical concepts, and has not lost any of its fascination for theorists up to date.

3.2.1

Fundamental Properties of Polymer Molecules

The characterizing property of polymers is their highly modular structure. They are composed of a large number of small building blocks (monomers), which are often all alike, but may also be combined to arbitrary sequences (in the case copolymers and biopolymers). The monomers are arranged in chains, which are usually flexible on the nanometer length scale, i.e., they can form kinks at little energetic expense, they curve around and may assume a large number of conformations at room temperature. The properties of such flexible polymers are largely determined by the entropy of the chain conformations. For example, the number of available conformations is reduced if molecules are stretched, which leads to a purely entropic restoring spring force [10] (rubber elasticity). Exposed to stress, polymeric systems respond by molecular rearrangements, which takes time and results in time-dependent strain (viscoelasticity). The fundamental processes that govern the behavior of polymeric materials do not depend on the chemical details of the monomer structure. For qualitative purposes, polymer molecules can be characterized by a few properties such as:

r r r r

The architecture of the molecules (linear chains, rings, stars, etc.); Physical properties (local chain stiffness, chain size, monomer volume); Physiochemical properties (monomer sequence, compatibility, charges); Special properties (e.g., a propensity to develop crystalline or liquid crystalline order).

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3.2.2

33

Coarse-Graining, Part I

The notion of ‘coarse-graining’ has lately become a buzzword in materials science, but the underlying concept is actually quite old in polymer science. The need for coarse-graining results from the fact that polymeric materials exhibit structure on very different length scales, ranging from Angstrom (the monomeric scale) to hundreds of nanometers (typical molecule extensions) or micrometers (supramolecular aggregates). It is not possible to treat all of them within one common theoretical framework. Therefore, different theoretical descriptions have been developed that deal with phenomena on different length and time scales. On the microscale, chemical details are taken into account and the polymers are treated at an atomistic level. This is the realm of theoretical chemistry. On the mesoscale, simplified molecule models come into play (string models, lattice models, bead-spring models, see below), whose behavior can be understood with concepts from statistical physics. Finally, on the macroscale, polymeric materials are described by continuous fields (composition, strain, stress, etc.) with certain mechanical properties, and their behavior can be calculated with methods borrowed from the engineers. In the following, we shall mainly focus on the mesoscale level, where polymers are described by extended molecules made of simplified ‘monomeric’ units, each representing several real monomers. Even within that level, one still has some freedom regarding the choice of the coarse-grained units. This is illustrated in Figure 3.1, where different coarse-grained representations of a polymer are superimposed onto each other. Polymers have remarkable universality properties, which allow one to link different coarse-grained representations in a rather well-defined manner, as long as the length scales under consideration are much larger than the (chemical) monomer length scale. For example, starting from one (atomistic or coarse-grained) model, we can construct a coarse-grained model by combining m ‘old’ units to one ‘new’ unit. If m is sufficiently large, the average squared distance d2  between two adjacent new units will depend on m according to a

Figure 3.1 Mesoscopic coarse-grained representations of a polymer molecule. Light pearl necklace in the background: Bead chain with bonds of fixed length. Solid and dashed lines: chain of coarse-grained units linked by bonds of variable length.

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characteristic power law d 2  ∼ m 2ν

(3.1)

where the exponent ν depends on the environment of a polymer, but not on chemical details [4, 5]. In a dense polymer melt, one has ν = 1/2 (see below). Similar scaling laws can be established for other chain parameters. 3.2.3

Ideal Chains

In mesoscopic polymer theories, one often uses as a starting point a virtual polymer chain where monomers that are well separated along the polymer backbone do not interact with each other, even if their spatial distance is small. Such polymers are called ‘ideal chains’. Even though they are mere theoretical constructions, they provide a good approximative description of polymers in melts and in certain solvents (‘Theta’-solvents, see below). 3.2.3.1

A Paradigm of Polymer Theory: The Gaussian Chain

Let us now consider a flexible ideal chain with N monomers, which we coarse-grain several times as sketched in Figure 3.1, until one coarse-grained monomer unites m ‘real’ monomers. For large m, the resulting chain is a random walk in space consisting of uncorrelated random steps di of varying length. According to the central limit theorem of probability theory [11], the steps are approximately Gaussian distributed, P(d) ∼ exp(−d 2 /2mσ 2 ), where σ does not depend on m. The same random walk statistics can be reproduced by a Boltzmann distribution with an effective coarsegrained Hamiltonian N /m 1 kB T  2 Hm = d 2 mσ 2 i=1 i

(3.2)

The Hamiltonian Hm describes the energy of a chain of springs with spring constant k B T m/σ 2 . The coarsegraining procedure has thus eliminated the information on chemical details (they are now incorporated in the single parameter σ ), and instead unearthed the entropically induced elastic behavior of the chain which lies at the heart of rubber elasticity. Eq. (3.2) is also an example for universal behavior in a polymer system (see section 3.2.2). The coarse-grained chain is self-similar. Every choice of m produces an equivalent model, provided the spring constant is rescaled accordingly. The distance between two coarse-grained units exhibits a scaling law of the form (3.1) as a function of m, d 2  = 3σ 2 m 2ν with ν = 1/2. Based on these considerations, it seems natural to define a ‘generic’ ideal chain model based on Eq. (3.2) with m = 1, N −1 1  HG [ri ] = (ri+1 − ri )2 kB T 2σ 2 i=1

(3.3)

the so-called ‘discrete Gaussian chain’ model. For theoretical purposes, it is often convenient to take the continuum limit: The index i in Eq. (3.2), which counts the monomers along the chain backbone, is replaced by a continuous variable s, the chain is parametrized by a continuous path R(s), and the steps d correspond

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35

to the local derivatives dR/ds of this path. The effective Hamiltonian then reads HG [R] 3 = 2 kB T 2b





N

ds 0

dR ds

2 (3.4)

This defines the continuous Gaussian chain. The only material parameters in Eq. (3.4) are the chain length N and the ‘statistical segment length’ or ‘Kuhn length’ b. Even those two are not independent, since they both depend on the definition of the monomer unit. An equivalent chain model can be obtained by rescaling N → N/λ and b2 → b2 λ. Hence the only true independent parameter is the extension of the chain, which can be characterized by the squared gyration radius Rg2

1 = N



N

ds (R(s) − R)2 = b2 N /6

(3.5)

0

where R = 1/N ∫ dsR(s) is the center of mass of the chain. The quantity Rg sets the (only) characteristic length scale of the Gaussian chain, and all length-dependent quantities scale with Rg . For example, the structure factor is given by  2  

1  N ds eik R(s)  = N g D k 2 Rg2 S(k) = N  0

(3.6)

with the Debye function g D (x) =

2 −x (e − 1 + x) x2

(3.7)

The Gaussian chain is not only a prototype model for ideal chains, it also provides a general framework for mesoscopic theories of polymer systems. The Hamiltonian HG (Eq. (3.4)) is then supplemented by additional terms that account for interactions, external fields, constraints (e.g., chemical crosslinks) etc. In this more general context, the Hamiltonian (3.4) is often referred to as ‘Edwards Hamiltonian’. Finally in this section, we note that from a mathematical point of view, the probability distribution of chain conformations defined by Eq. (3.4) is a Wiener measure [11, 12]. The continuum limit leading to Eq. (3.4) is far from trivial, but well-defined. We shall not delve further into this matter. 3.2.3.2

Other Chain Models

The Gaussian chain model is a common starting point for analytical theories of long flexible polymers on sufficiently large length scales. On smaller length scales, or for stiffer polymers, or for computer simulation purposes, other types of coarse-grained models have proven useful. We briefly summarize some popular examples. The wormlike chain model is a continuous model designed to describe stiff polymers. They are represented by smooth paths R(s) with fixed contour length N, where the parameter s runs over the arc length of the curve, i.e., the derivative vector u = dR/ds has length unity, |u| ≡ 1. The paths have a bending stiffness η, such that they are distributed according to the effective Hamiltonian η HW LC [R] = kB T 2





N

ds 0

d2 R ds 2

2 (3.8)

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The wormlike chain model is particularly useful if local orientational degrees of freedom are important. The freely jointed chain is a discrete chain model where the chain is composed of N links of fixed length. It is often used to study general properties of ideal chains. The spring-bead chain is a chain of beads connected with springs. It has some resemblance with the discrete Gaussian chain model, except that the springs have a finite equilibrium length. Spring-bead models are popular in computer simulations. In lattice models, the monomer positions are confined to the sites of a lattice. This simplifies both theoretical considerations and computer simulations. 3.2.4

Interacting Chains

The statistical properties of chains change fundamentally if monomers interact with each other. Such interactions are readily introduced in the coarse-grained models presented above. In the discrete models, one simply adds explicit interactions between monomers. In the continuous path models, one supplements the energy contribution for individual ideal chains, Eq. (3.4) or (3.8), by an interaction term, such as v ˆ = H I [ρ] 2



w dr ρˆ + 6 2

 dr ρˆ 3 + · · ·

(3.9)

(for weak interactions), where the ‘monomer density’ ρ(r ˆ ) is defined as ρ(r) ˆ =

 α

N

ds δ(r − Rα (s))

(3.10)

0

and the sum α runs over the polymers Rα (s) in the system. Eq. (3.9) corresponds to a virial expansion of the local interaction energy in powers of the density. In many cases, only the quadratic term (v) needs to be taken into account (‘two-parameter Edwards model’). The Ansatz (3.9) is suitable for dilute polymer systems – dense systems are discussed below (Section 3.2.4.2). The generalization to multiphase systems where monomers may have different type A, B, . . . is straightforward. One simply operates with different densities ρˆ A , ρˆ B , . . . and interaction parameters v A A , v AB , . . .. Interactions complicate the theoretical treatment considerably and in general, exact analytical solutions are no longer available. The properties of interacting polymer systems have been explored theoretically within mean-field approximations, renormalization-group calculations, scaling arguments, and computer simulations. To set the stage for the discussion of multiphase systems in Sections 3.3 and 3.4, we will now briefly sketch the most important scenarios for monophase polymer systems. 3.2.4.1

Polymers in Solution and Blobs

We first consider single, isolated polymer chains in solution. Their properties depend on the quality of the solvent, which is incorporated in the second virial parameter v in Eq. (3.9) (the three-body parameter w is typically positive [13]). In good solvent (v > 0), monomers effectively repel each other, and the chain swells. Extensive theoretical work [4] has shown that the scaling behavior (Eq. (3.1)) remains valid, but the exponent ν increases from ν = 1/2 for ideal chains to ν ≈ 3/(d + 2), where d is the spatial dimension (more precisely, ν = 0.588 in three dimensions). This is the famous ‘Flory exponent’, which characterizes the scaling behavior of so-called ‘self-avoiding chains’. Accordingly, the gyration radius of the chain scales

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with the chain length like Rg2 ∼ N 2ν

(3.11)

In bad solvent (v < 0), monomers effectively attract each other and the chain collapses. At the transition between the two regimes, the ‘Theta point’ (v θ ≈ 0), the scaling behavior basically corresponds to that of ideal chains (ν = 1/2), except for subtle corrections due to the three-body w-term [4]. Eq. (3.11) describes the behavior of single, unperturbed chains. Even in good solvent, the self-avoiding scaling is often disturbed. For example, the chains cannot swell freely if they are confined, or if they are subject to external forces. Another important factor is the concentration of chains in the solution: If many chains overlap, the intrachain interactions are screened on large length scales. Loosely speaking, monomers cannot distinguish between interactions with monomers from the same chain and from other chains. As a result, chains no longer swell and ideal chain behavior is recovered. This mechanism applies in three or more spatial dimensions. Two dimensional chains segregate [4, 14, 15]. All of these situations can be analyzed within one single ingenious framework, the ‘blob’ picture introduced Daoud et al. in 1975 [17]. It is based on the assumption that there exists a crossover length scale ξ below which the chain is unperturbed. Blobs are volume elements of size ξ within which the polymers behave like self-avoiding chains. On larger scales, the polymer behaves like an ideal chain consisting of a string of blobs. Every blob contains m ∼ ξ 1/ν monomers and carries a free energy of the order kB T. These simple rules are the whole essence of the blob model. We shall illustrate their use by applying them to a number of prototype situations depicted in Figure 3.2. Concentrated polymer solution (Figure 3.2 a). For polymer concentrations , we calculate the crossover length scale ξ from self-avoiding to ideal behavior. Since ξ is the blob size, we can simply equate = m/ξ 3 , i.e., ξ ∼ φ −ν/(3ν−1) . Polymer confined in a slit (Figure 3.2 b). We consider the free energy penalty F on the confinement. Here, the blob size is set by the width R of the slit. Each blob contains m ∼ R1/ν monomers, hence the total free energy scales like F ∼ N/m ∼ NR−1/ν . Polymer confined in a cavity (Figure 3.2 c) The result b) also holds for chains confined in a tube. In closed cavities, however, the situation is different due to the fact that the cavity constrains the monomer concentration. The resulting blob size is ξ ∼ (N/R3 )ν /(1−3ν) , and the free energy of confinement scales as (a)

(b)

(c)

(d)

Figure 3.2 Illustration of the blob model in different situations: (a) concentrated polymer solution; (b) Polymer confined in a slit; (c) chain confined in a spherical cavity; (d) structure formation in solutions of miktoarm star copolymers (after Ref. [16]). See text for explanation.

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F ∼ (R/ξ )3 ∼ (R/N ν )3/(1−3ν) . This has been discussed controversially, but was recently confirmed by careful computer simulations [18]. ABC miktoarm star copolymers in selective solvent (Figure 3.2d). Last, we cite a recent application to a multiphase polymer system. Zhulina and Borisov [16] have studied ABC star copolymers by means of scaling arguments. They derived a rich state diagram, according to which ABC star copolymers may assemble to several types of nanostructures, among other spherical micelles, dumbbell micelles, and striped cylindrical micelles. This is only one of numerous examples where scaling arguments have been used to analyze complex multicomponent systems.

3.2.4.2

Dense Melts

Dense melts can be considered as extreme cases of a very concentrated polymer solution, hence it is not surprising that the chains effectively exhibit ideal chain behavior. In fact, the situation is more complicated than this simple argument suggests. The quasi-ideal behavior results from a cancellation of two effects: On the one hand, the intrachain interactions promote chain swelling, but on the other hand, the chain pushes other chains aside (‘correlation hole’), which in turn exert pressure and squeeze it. Deviations from true ideal behavior can be observed, e.g., at the level of chain orientational correlations [19, 20]. Nevertheless, the ideality assumption is a good working hypothesis in dense melts and shall also be used here in the following.

3.2.5

Chain Dynamics

In this chapter we focus on equilibrium and static properties of polymer systems. We can only touch on the possible dynamical behavior, which is even more diverse. In the time scales of interest, the motion of polymers is diffusive, i.e., the inertia of the macromolecules is not important. Three prominent types of dynamical behavior have been established. In the Rouse regime, the chain dynamics is mainly driven by direct intrachain interactions. This regime is encountered for short chains. The dynamical properties of ideal chains can be calculated exactly, and the results can be generalized to self-avoiding chains using scaling arguments. One of the important properties of Rouse chains is that their sedimentation mobility does not depend on the chain length N. Hence the diffusion constant scales like D ∼ 1/N , and the longest internal relaxation time, which can be estimated as the time in which the chain diffuses a distance Rg , scales like τ ∼ N 2ν+1 . In the Zimm regime, the dynamics is governed by long-range hydrodynamic interactions between monomers. This regime develops for sufficiently long chains in dilute solution. They diffuse like Stokes spheres with the diffusion constant D ∼ 1/Rg , and the longest relaxation time scales like τ ∼ Rg3 . In concentrated solutions, the hydrodynamic interactions are screened [5] and Rouse behavior is recovered after an initial Zimm period [21]. The reptation regime is encountered in dense systems of chains with very high molecular weight. In this case, the chain motion is topologically constrained by the surrounding polymer network, and they are effectively confined to move along a tube in a a snake-like fashion [4, 22]. The diffusion constant of linear chains scales like D ∼ 1/N 2 and the longest relaxation time like τ ∼ N 3 . This description is very schematic and oversimplifies the situation even for fluids of linear polymers. Moreover, most polymer materials are not in a pure fluid state. They are often cooled down below the glass transition, or they partly crystallize – in both cases, the dynamics is frozen. Chemical or physical crosslinks constrain the motion of the chains and impart solid-like behavior. In multiphase polymer systems, the situation is further complicated by the fact that the glass point or the crystallization temperature of the

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Spheres (S)

Cylinders (C)

Gyroid (G)

39

Lamellae (L)

Figure 3.3 Self-assembled copolymer mesophases.

different components may differ, such that one component freezes where the other still remains fluid [23–32]. In the following discussion, we shall limit ourselves to fluid multiphase polymer systems.

3.3 Theory of Multiphase Polymer Mixtures After this general overview, we turn to the discussion of polymer blends. We consider dense mixtures, where the polymers are in the melt regime (Section 3.2.4.2). Moreover, we assume incompressibility – the characteristic length scales of density fluctuations are taken to be much smaller than the length scales of interest here. Monomers of different type are usually slightly incompatible (see Section 3.3.1.3). In polymers, the incompatibilities are amplified, such that macromolecules of different type tend to be immiscible: Blended together, they demix and develop an inhomogeneous multiphase structure where microdroplets of one phase are finely dispersed in another phase. In order to overcome or at least control this situation, copolymer molecules can be added in which the two incompatible components are chemically linked to each other. They act as compatibilizer, i.e., they shift the demixing transition and reduce the interfacial tension between different phases in the demixed region. At high concentrations, they are found to self-organize into a variety of ordered mesophases (microphase separation; see, e.g., structures shown in Figure 3.3). Hence copolymers can also be used to manufacture nanostructured materials in a controlled way. Nowadays, the theory of structure formation in polymer blends has reached a highly advanced level and theoretical calculations have predictive power, e.g., with respect to structures that can be expected in new polymeric materials. In this section, we shall present some of the most successful theoretical approaches. 3.3.1

Flory-Huggins Theory

We begin with sketching the Flory-Huggins theory, which is the classical theory of phase separation in polymer blends, and which in some sense lays the foundations for all later, more sophisticated theories of polymer mixtures. 3.3.1.1

Basic Model for Binary Blends

We consider a binary blend of homopolymers A and B with length NA and NB , and volume fractions A and B . According to Flory [1] and Huggins [2], the free energy per monomer is approximately given by fF H B A ln( A ) + ln( B ) + χ A B = kB T NA NB

(3.12)

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4 N =NA/2 B

3

χNA

3

demixed

χNA

2

N =NA B

2 N =2NA

1

B

N =2N B

0 0

A

1

mixed

ΦA

1

0

ΦA

0 1

Figure 3.4 Left: Phase diagram for a binary AB polymer blend with B-chains twice as long as A-chains according to the Flory-Huggins theory. Thick solid line shows the binodal line (i.e., the demixing line), thin dashed line the spinodal line (i.e., the line where the homogeneous blend becomes unstable). Right: Binodals for binary AB blends with different chain length ratios as indicated.

with A + B = 1. The first two terms account for the mixing entropy of the two components, and the last term for the (in)compatibility of the monomers. The parameter χ is the famous ‘Flory Huggins parameter’, which will be discussed in more detail below. The generalization of this expression to ternary ABC homopolymer blends, etc. is straighforward, one only needs to introduce several χ -parameters χ AB , χ BC , and χ AC . Here we will only discuss binary systems. By minimizing the free energy, Eq. (3.12), one easily identifies the region in phase space where the mixture phase separates into an A-rich phase and a B-rich phase. At low of χ , the blend remains √ √ values 2 = (1/ N + 1/ N ) for the critical composition homogeneous. Demixing sets in at a critical value 2χ c A B √ A,c = 1/(1 + N A /N B ). The region of stability of the homogeneous (mixed) blend is delimited by the ‘binodal’ line (see Figure 3.4). Beyond the binodal, the homogeneous blend may still remain metastable. It becomes unstable at the ‘spinodal’, which is defined as the line where the second derivative of fFH in Eq. (3.12) with respect to A vanishes. An example of a phase diagram with a binodal and a spinodal is shown in Figure 3.4 (left). Figure 3.4 (right) demonstrates the shift of the binodal with varying chain length ratio NA /NB . The Flory-Huggins free energy, Eq. (3.12), was originally derived based on a lattice model, but it can also be applied to off-lattice systems. It does, however, rely on three critical assumptions:

r r r

The polymer conformations are taken to be those of ideal chains, independent of the composition (ideality assumption, cf. Section 3.2.4.2). The melt is taken to be incompressible, and monomers A and B occupy equal volumes. Local composition fluctuations are neglected (mean-field assumption).

In reality, none of these assumptions is strictly valid. The polymer conformations do depend on the composition, most notably for chains of the minority component. The incompressibility assumption is reasonable, but the volumes per monomer are not equal. As a consequence, the χ -parameter is not a fixed parameter (at fixed temperature), but depends on the composition of the blend (see Section 3.3.1.3). Finally, the composition fluctuations shift phase boundaries and may even fundamentally change the phase behavior. (see Section 3.3.2.5).

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3.3.1.2

41

Inhomogeneous Systems: Flory-Huggins-de Gennes Theory

Equation (3.12), only describes homogeneous systems. The simplest approach to generalizing the FloryHuggins theory to inhomogeneous systems, e.g., polymer blends containing interfaces, consists in adding a penalty on composition variations (∇ A )2 = (∇ B )2 . The coefficient of the square gradient term can be derived within a more advanced mean-field treatment, the random phase approximation, which will be described further below (Section 3.3.3.1). One obtains the Flory-Huggins-de Gennes free energy functional for polymer blends, 

 FFHdG [ A (r)] = ρ0

dr

f F H ( A (r)) +

kB T 36



b2A b2 + B A B



(∇ A )2

(3.13)

(with B = 1 − A ), where bA and bB are the Kuhn lengths of the homopolymers A and B, and fFH ( A ) is given by Eq. (3.12). The functional (3.13) can be applied if composition variations are weak, and have characteristic length scales of the order of the gyration radius of the chains (‘weak segregation regime’, see Section 3.3.3.1). A very similar functional can be derived in the opposite case, where A- and B-polymers are fully demixed and separated by narrow interfaces. In this ‘strong segregation’ regime, the blend can be described by the functional (see Section 3.3.3.2)  FSSL [ A (r)] = ρ0



kB T dr χ A B + 24



b2A b2 + B A B



(∇ A )

2

(3.14)

We note that at strong segregation, the mixing entropy terms in in fFH , Eq. (3.12), can be neglected, hence the functionals (3.13) and (3.14) are identical except for the numerical prefactor of the square gradient term. In the strong segregation limit, the square gradient penalty results from an entropic penalty on A and B segments due to the presence of the interface, whereas in the weak segregation limit, it is caused by the deformation of whole chains. 3.3.1.3

Connection to Reality: The Flory-Huggins Parameter

In the Flory-Huggins theory, the microscopic features of the blend are incorporated in the single Flory-Huggins parameter χ . Not surprisingly, this parameter is very hard to access from first principles. In the original Flory-Huggins lattice model, χ is derived from the energetic interactions between monomers that are neighbors on the lattice. The interaction energy between monomers i and j is taken to be characterized by energy parameters  ij . The χ -parameter is then given by χ=

z−2 (2 AB −  A A −  B B ) 2k B T

(3.15)

where z is the coordination number of the lattice. It is reduced by two (z − 2) in order to account for the fact that interactions between neighbor monomers on the chain are fixed and have no influence on the demixing behavior. In reality, the situation is not as simple. The miscibility patterns in real blends tend to deviate dramatically from that predicted by the Flory-Huggins model. Several blends exhibit a lower critical point instead of an upper critical point, indicating that the demixing is driven by entropy rather than enthalpy. The critical temperature Tc of demixing often does not scale linearly with the chain length N, as one would expect from

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Eq. (3.15) (see Figure 3.4). In blends of polystyrene and poly(vinyl methyl ether) (PS/PVME), for example, Tc is nearly independent of N, and the critical concentration c is highly asymmetric even for roughly equal chain lengths, in apparent contrast with Figure 3.4 (NB = NA ) [33]. Formally, this problem can be resolved by arguing that the Flory-Huggins expression for χ , Eq. (3.15), is oversimplified and that χ kB T is not a constant. Several factors contribute to the χ -parameter, leading to a complex dependence on the temperature, the blend composition, and even the chain length. Monomer incompatibility: Monomers may be incompatible both for enthalpic and entropic reasons. For example, consider two nonpolar monomers i and j. The van-der Waals attraction between them is proportional to the product of their polarizabilities α, hence  ij ∝ α i α j and χ ∝ (α A − α B )2 . More generally, the enthalpic incompatibility of monomers i and j can be estimated by χ H ∝ (δ A − δ B )2 , where the δ i are the Hildebrand solubility parameters of the components [34, 35]. In addition, entropic factors may contribute to the monomer incompatibility, which are, e.g., related to shape or stiffness disparities [36–40]. Since the enthalpic and entropic contributions evolve differently as a function of the temperature, the χ -parameter will in general exhibit a complicated temperature dependence. Equation-of-state effect: In general, the volume per polymer depends on the composition of the blend. Already the volume per monomer is usually different for different monomer species; at constant pressure, χ therefore varies roughly linearly with A [41]. In blends of monomers with very similar monomer structure, e.g., isotopic blends, the linear contribution vanishes and a weak parabolic dependence remains, which can partly [42] (but not fully [43]) be explained by an ‘excess volume of mixing’. Chain correlations: Since the demixing is driven by intermolecular contacts, intramolecular contacts only contribute indirectly to the χ -parameter. The estimate for χ can be improved if one replaces the factor z − 2 by an ‘effective coordination number’ zeff which is given by the mean number of interchain contacts per monomer [44]. Moreover, the ideality assumption (see Section 3.2.4.2) is not strictly valid. Chains in the minority phase (e.g., A chains in a B-rich phase) tend to shrink in order to reduce unfavorable contacts [45]. Since the segregation is effectively driven by χ N, χ slightly depends on the chain length N as a result. The situation becomes even more complex if copolymers are involved, which assume dumbbell shapes even in a disordered environment [46, 47]. This also may affect the effective χ -parameter [48]. Composition correlations: The effective interactions between monomers change if the local environment is not random. We have already noted earlier that the composition may fluctuate. Large-scale fluctuations can be incorporated in the Flory-Huggins framework in terms of a fluctuating field theory (see Sections 3.3.2.5 and 3.3.3.1). Fluctuations (correlations) on the monomeric scale renormalize the χ -parameter (nonrandom mixing). Moreover, the local fluid structure may depend on the local composition (nonrandom packing). In view of these complications, establishing an exhaustive theory of the χ -parameter remains a formidable task [37, 49–54]. Even the reverse problem of designing simplified particle-based polymer models with a well-defined χ -parameters turns out to be highly non-trivial [55]. The very concept of a χ -parameter has been challenged repeatedly, e.g., by Tambasco et al. [56], who analyzed experimental data for a series of blends and found that their thermodynamic behavior can be related to a single ‘g − 1-parameter’, which is independent of composition, temperature and pressure. They suggest that this parameter may be more appropriate to characterize blends than the χ -parameter. However, it has to be used in conjunction with an integral equation theory, the BGY lattice theory by Lipson [57–59], which is much more involved than the Flory-Huggins theory especially when applied to inhomogeneous systems. Freed and coworkers [36, 37, 40, 60] have proposed a generalized Flory-Huggins theory, the ‘lattice cluster theory’, which provides a consistent microscopic theory for macroscopic thermodynamic behavior. In a certain limit (high pressure, high molecular weight, fully flexible chain), this theory reproduces a Flory-Huggins type

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free energy with an effective χ -parameter [49] χeff

  2( p A A + p B B ) (r A − r B )2 = + χ0 1 − z2 z(z − 2)

(3.16)

where χ 0 is given by Eq. (3.15) and the parameters rj , pj depend on the structure of the monomers j. Dudowicz et al. have recently pointed out that this model can account for a wide range of experimentally observed miscibility behavior [38, 39]. Another pragmatic and reasonably successful approach consists in using χ as a heuristic parameter. Assuming that it is at least independent of the chain length, it can be determined experimentally, e.g., from fitting small-angle scattering data to theoretically predicted structure factors [61–63]. Alternatively, χ can be estimated from atomistic simulations [55, 64]. The results from experiment and simulation tend to compare favorably even for systems of complex polymers [65]. 3.3.2

Self-consistent Field Theory

We have taken some care to discuss the Flory-Huggins theory because it establishes a framework for more general theories of polymer blends. In particular, it provides the concept of the Flory-Huggins parameter χ , which we will take for granted from now on, and take to be independent of the composition, despite the question marks raised in the previous section (Section 3.3.1.3). In this section, we present a more sophisticated mean-field approach, the self-consistent field (SCF) theory. It was first proposed by Helfand and coworkers [66–70] and has since evolved to be one of the most powerful tools in polymer theory. Reviews on the SCF approach can be found, e.g., in [8, 71, 72]. 3.3.2.1

How It Works in Principle

For simplicity, we will first present the SCF formalism for binary blends, and discuss possible extensions later. Our starting point is the Edwards Hamiltonian for Gaussian chains, Eq. (3.4), with a Flory-Huggins interaction term,  H I [ρˆ A , ρˆ B ]/k B T = ρ0 χ

ˆ A ˆB dr

(3.17)

ˆ j = ρˆ j /ρ0 and incompressibility is requested, i.e., where we have defined the ‘monomer volume fractions’ ˆ B ≡ 1 everywhere. The quantity ρˆ j depends on the paths Rα of chains of type j and has been defined ˆ A+ in Eq. (3.10). We consider a mixture of nA homopolymers A of length NA and nB homopolymers B of length NB in the volume V. The canonical partition function is given by 

   1 ˆ ˆ −HG [Rα ]/k BT ˆ A+ ˆ B − 1) Z= ρ0 DRα e e−ρ0 χ dr A B δ( n A !n B ! α

(3.18)

 The product α runs over all chains in the system, ρ 0 DRα denotes the path integral over all paths Rα (s), and we have introduced the factor ρ 0 in order to make Z dimensionless. The path integrals can be decoupled by inserting delta-functions ∫ Dρ j δ(ρ j − ρˆ j ) (with j = A,B), 1 and using the Fourier representations of the delta-functions δ(ρ j − ρˆ j ) = ∫i∞ DW j e N W j (ρ j −ρˆ j ) and

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δ(ρ A + ρ B − ρ 0 ) = ∫i∞ Dξ e N ξ (ρ A +ρ B −ρ0 ) . Here N is some reference chain length, and the factor 1/N is introduced for convenience. This allows one to rewrite the partition function in the following form: 1



 DW A DW B Dξ

Z∝ i∞

∞

Dρ A Dρ B e−F/k B T

(3.19)

with ⎧    F[WA , WB , ξ, ρA , ρB ] ρ0 ⎨ dr W j j = χ N dr A B − kB T N ⎩ j=A,B

 −

⎫ ⎬  Qj N dr ξ ( A + B − 1) − Vj ln ρ0 N nj ⎭ j j=A,B 

(3.20)

( j = ρ j /ρ 0 ), where Vj denotes the partial volume occupied by all polymers of type j in the system, and the functional   Nj 1 (3.21) Q j = DR e−HG [R]/k B T e− N 0 ds W j (R(s)) is the partition function of a single, noninteracting chain j in the external field Wj . Thus the path integrals are decoupled as intended, and the coupling is transferred to the integral over fluctuating fields Wj and ξ . Now the self-consistent field approximation consists in replacing the integral (3.19) by its saddle point, i.e., minimizing the effective Hamiltonian H with respect to the variables ρ j (r) and Wj (r). The minimization procedure results in a set of equations, ˆ j = j  ˆ i − ξ W j = χ N 

with with

j = A, B i, j = A, B

and i = j

(3.22)

ˆ j  denotes the average of ˆ j in a system of noninteracting chains subject to the external fields where  Wj (r). We note that the latter are real, according to Eq. (3.22), even though the original integral (3.19) is carried out over the imaginary axis. Intuitively, the Wj (r) can be interpreted as the effective mean fields acting on monomers due to the interactions with the surrounding monomers. Together with the incompressibility constraint, the equations (3.22) form a closed cycle which can be solved self-consistently. For future reference, we note that it is sometimes convenient to carry out the saddle point integral only with respect to the variables Wj (r). This defines a free energy functional FSCF [ A ], which has essentially the same form as F (Eq. 3.20), except that the variables W α (r) and ξ (r) are now real Lagrange parameter fields ˆ B  = B , and B = 1 − A and depend on A . The same functional can also ˆ A  = A ,  that enforce  be derived by standard density functional approaches, using as the reference system a gas of non-interacting Gaussian chains [73]. In some cases, one would prefer to operate in the grand canonical ensemble, i.e., at variable polymer numbers nj . The resulting SCF theory is very similar. The last term in Eq. (3.25) is replaced by (− j z j Q j ), where zj is proportional to the fugacity of the polymers j. The formalism can easily be generalized to other inhomogeneous polymer systems. The application of the theory to multicomponent A/B/C/. . . homopolymer blends with a more general interaction Hamiltonian H I [ρˆ A , ρˆ B , ρˆC , . . .] replacing Eq. (3.17) is straightforward. The self-consistent field equations (3.22) are

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simply replaced by

Wj =

N δH I [{ρi }] −ξ kB T δρ j

(3.23)

If copolymers are involved, the single-chain partition function Qc for the corresponding molecules must be adjusted accordingly. For example, the single-chain partition function for an A:B diblock copolymer of length Nc with A-fraction f reads  Qc =

1

DR e−HG [R]/k B T e N



f Nc 0

ds W A (R(s))+ N1

 Nc f Nc

ds W B (R(s))

(3.24)

The general SCF free energy functional for incompressible multicomponent systems is given by     FSCF [{ρ j }] H I [{ρi }]  1 1 = − ρ j − ρ0 dr W j ρ j − dr ξ kB T kB T N N j j    Qα − n α ln ρ0 nα α

(3.25)

where the sum j runs over monomer species and the sum α over polymer types. The SCF theory has been extended in various ways to treat more complex systems, e.g., compressible melts and solutions [74, 75], macromolecules with complex architecture [76], semiflexible polymers [77] with orientational interactions [78, 79], charged polymers [80], polydisperse systems [81, 82], polymer systems subject to stresses [83, 84], systems of polymers undergoing reversible bonds [85, 86], or polymer/colloid composites [87–89].

3.3.2.2

How It Works in Practice

In the previous section, we have derived the basic equations of the SCF theory. Now we describe how to solve them in practice. The first task is to evaluate the single-chain partition functions Q j and the corresponding density averages ρˆ j  for noninteracting chains in an external field. We consider a single ideal chain of length N in an external field W(r, s), which may not only vary in space r, but also depend on the monomer position s in the chain in the case of copolymers. It is convenient to introduce partial partition functions  q(r, s) = q † (r, s) =



DR e−HG [R]/k B T e N 1

DR e−HG [R]/k B T e N 1

s 0

s 0

ds  W (R(s  ),s  )

δ(R(s) − r)

ds  W (R(s  ),N −s  )

δ(R(s) − r)

(3.26) (3.27)

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where the path integrals DR are carried out over paths of length s. According to the Feynman-Kac formula [90], the functions q(r, s) and q† (r, s) satisfy a diffusion equation  2  b W (r, s) ∂ q(r, s) = − q(r, s) ∂s 6 N   2 W (r, N − s) ∂ † b q (r, s) = − q † (r, s) ∂s 6 N

(3.28) (3.29)

with the initial condition q(r, 0) = q† (r, 0) = 1. Numerical methods to solve diffusion equations are available [91], hence q and q† are accessible quantities. The single-chain partition function can then be calculated via   Q = dr q(r, N ) = dr q † (r, N ) (3.30) and the distribution of the sth monomer in space is q(r, s) q† (r, N − s)/Q. More specifically, to study binary homopolymer blends, one must solve the diffusion equations for the † partial partition functions q j = q j in an external field W = Wj (with j = A, B). The averaged volume fractions of monomers j are then given by ˆ j (r) = 

1 nj ρ0 Q j

 0

Nj

†

ds q j (r, s) q j (r, N − s)

(3.31)

in the canonical ensemble, and ˆ j (r) = z j  N



Nj 0

†

ds q j (r, s) q j (r, N − s)

(3.32)

in the grand canonical ensemble. If AB diblock copolymers with fraction f of Amonomers are present, one † must calculate the partial partition functions qc and qc in the external field W(r, s) = WA for s < f Nc and ˆ A  is n C /ρ0 QC f A (r) W(r) = WB for s ≥ f Nc . The contribution of the copolymers to the volume fraction   fN † with f A (r) = 0 c ds qC (r, s)qC (r, N − s) in the canonical ensemble, and z c /N fA (r) in the grandcanonical ensemble. With this recipe at hand, one can calculate the different terms in Eqs. (3.22). The next problem is to solve these equations simultaneously, taking account of the incompressibility constraint. This is usually done iteratively. We refer the reader to Section 3.4. in [91] for a discussion of different iteration methods. 3.3.2.3

Application: Diblock Copolymer Blends, Part I

To illustrate the power of the SCF approach, we cite one of its most spectacular successes: The reproduction of arbitrarily complex copolymer mesophases. In a series of seminal papers, Matsen and coworkers have calculated phase diagrams for diblock copolymer melts [95, 96]. Figure 3.5 compares an experimental phase diagram due to Bates and coworkers [92–94] with the SCF phase diagram of Matsen and Bates [95]. The SCF theory reproduces the experimentally observed structures. At high values of χ N (‘strong segregation’) the SCF phase diagram features the correct sequence of mesophases at almost the correct value of the fraction of A-monomers f . At low values of χ N (‘weak segregation’), the two phase diagrams are distinctly different.

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47

40

C

S

C

L

30

30

G

χ N 20

χN 20 10

S

L

C

Scp

C

S Scp

G

10 disordered

0 0.0

0.2

0.4

0.6

disordered 0.8

f

1.0

0 0.0

0.2

0.4

0.6

0.8

1.0

f

Figure 3.5 Experimental phase diagram for polystyrene-polyisoprene diblock copolymer melts (left, after Refs. [92–94]) compared with phase diagram obtained with the SCF theory by Matsen and coworkers (right, after Ref. [95]) in the coordinates χ N and A block volume fraction f. The labels S,C,L, and G correspond to the structures shown in Figure 3.3. In addition, the SCF phase diagram features a close-packed sphere phase Scp . All phase transitions are first order except for the disordered/lamellar-transition at f = 1/2 in the SCF phase diagram. Courtesy of Mark W. Matsen, adapted from M.W. Matsen, J. Phys.: Cond. Matter 14, R21 (2002).

This can be explained by the effect of fluctuations and will be discussed further below (Sections 3.3.2.5 and 3.4.4.2, see also Figure 3.8). Tyler and Morse have recently reconsidered the SCF phase diagram and predicted the existence of yet another mesophase, which has an orthorombic unit cell and an Fddd structure and intrudes in a narrow regime at the low χ N-end of the gyroid phase [97]. This phase was later indeed found in a polystyrene-polyisoprene diblock copolymer melt by Takenaka and coworkers [98]. 3.3.2.4

Related Mean Field Approaches

So far, we have focussed on sketching a variant of the SCF theory which was originally developed by Helfand and coworkers [70]. A number of similar approaches have been proposed in the literature. Scheutjens and Fleer have developed a SCF theory for lattice models [99], which is applied very widely [6]. Scheutjens-Fleer calculations are very efficient and incorporate in a natural way the finite (nonzero) range of monomer interactions. To account for this in the Helfand theory, one must introduce additional terms in Eq. (3.9) [69, 70], which indeed turn out to become important in the vicinity of surfaces [75]. Carignano and Szleifer [100] have proposed an SCF theory where chains are sampled as a whole in a surrounding mean field. Hence intramolecular interactions are accounted for exactly and the chain statistics correspond to that of self-avoiding walks (Section 3.2.4.1). This approach is more suitable than the standard SCF theory to study polymers in solution, or melts of molecules with low-molecular weight, where the ideality assumption (see Section 3.2.4.2) becomes questionable [101]. In this chapter, we have chosen a field-theoretic way to present the SCF theory. Freed and coworkers [73, 102] have derived the same type of theory from a density functional approach, using a reference system of non-interacting Gaussian chains. Compared to the density functional approach, the field-theoretic approach has the advantage that the effect of fluctuations can be treated in a more transparent way (see Section 3.3.2.5). On the other hand, information on the local liquid structure of the melt (i.e., monomer correlation functions, packing effects. etc.), can be incorporated more easily in density functional approaches [103, 104]. Density functionals have also served as a starting point for the development of dynamical theories which allow to study the evolution of multiphase polymer blends in time [105–108] (see Section 3.4.3.2).

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10

Disordered

140 Disordered

III 11

T (°C)

100 2φ & 3φ

60

20

0

0.2

0.4

II

2φ

DμΕ

12

G μΕ IV

Lamellar

13

BμE –20

χΝ

0.6

φHomopolymer

0.8

1

0.4

Lamellar 0.5

0.6

0.7

0.8

0.9

1

φHomopolymer

Figure 3.6 Experimental phase diagram for symmetric ternary blends of PDMS+PEE+PDMS-PEE (NC ∼ 5NH ) featuring lamellar phase, phase-separated region 2 , and microemulsion channel BμE (left, from Ref. [109]), compared with theoretical phase diagrams (right) from SCF theory (solid lines) and Monte Carlo simulations at C = 50 (dashed lines/circles). The dotted lines separate disordered regions with different local structure: DμE, defect driven disorder, and GμE, genuine microemulsion morphology. The regions I–IV are discussed in the text. Left Figure: Reprinted from [454]. Copright (1999) with permission from Royal Society of Chemistry.

3.3.2.5

Fluctuation Effects

Mean-field approaches for polymer systems like the SCF theory tend to be quite successful, because polymers overlap strongly and have many interaction partners. However, there are several instances where composition fluctuations become important and may affect the phase behavior qualitatively. To illustrate some of them, we show the phase diagram of ternary mixtures containing A and B homopolymers and AB diblock copolymers in Figure 3.6. The left graph shows the experimental phase diagram [109], the right graph theoretical phase diagrams obtained by D¨ouchs et al. [110, 111] from the SCF theory (solid lines) and from field-based computer simulations (dashed line, see Section 3.4.3.2 for details on the simulation method). Regions where different types of fluctuations come into play are marked by I–IV.

I) Fluctuations are important in the close vicinity of critical points, i.e., continuous phase transitions. They affect the values of the critical exponents, which characterize, e.g., the behavior of the specific heat at the transition [90]. In Figure 3.6, such critical transitions are encountered at high homopolymer concentration, where the system essentially behaves like a binary A/B mixture with a critical demixing point. This point belongs to the Ising universality class, hence the system should exhibit Ising critical behavior. It has to be noted that in polymer blends, critical exponents typically remain mean-field like until very close to the critical point [112–114]. II) The effect of fluctuations is more dramatic in the vicinity of order-disorder transitions (ODT), e.g., the transition between the disordered phase and the lamellar phase at low homopolymer concentrations. Fluctuations destroy the long-range order in weakly segregated periodic structures, they shift the ODT and change the order of the transition from continuous to first order (Brazovskii mechanism [115, 116]).

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This effect accounts for the differences between the experimental and the SCF phase diagram in Figure 3.5. III) The SCF phase diagram features a three-phase (Lamellar + A + B) coexistence region reaching up to a Lifshitz point. Lifshitz points are generally believed to be destroyed by fluctuations in three dimensions. IV) In strongly segregated mixtures, fluctuations affect the large-scale structure of interfaces. Whereas mean-field interfaces are flat, real interfaces undulate. The so-called ‘capillary waves’ may destroy the orientational order in highly swollen lamellar phases. A locally segregated, but globally disordered ‘microemulsion’ state intrudes between the homopolymer-poor lamellar phase and the homopolymerrich two-phase region in Figure 3.6.

Both in the cases of III) and IV), the effect of fluctuations is to destroy lamellar order in favor of a disordered state. However, the mechanisms are different. This is found to leave a signature in the structure of the disordered phase, which is still locally structured with a characteristic wavevector q* [111]. In the Brazovskii regime, the wavevector q* corresponds to that calculated from the SCF theory (defect driven disorder regime, DμE). In the capillary wave regime, the characteristic length scale increases, compared to that calculated from the SCF theory (genuine microemulsion regime, GμE). Formally, the effect of fluctuations is hidden in the overall prefactor ρ 0 /N in the SCF Hamiltonian H (Eq. 3.25). The larger this factor, the more accurate is the saddle point integration that lies at the heart of the SCF approximation. We can thus define a ‘Ginzburg parameter’ C = Rgd ρ0 /N , which characterizes the strength of the fluctuations. Here the factor Rgd must be introduced to make C dimensionless (d is the spatial dimension), and Rg = Nb2 is the natural length scale of the system, the radius of gyration of an ideal chain of length N. The Ginzburg parameter roughly corresponds to the ratio of the volume spanned by a chain, Rgd , and the volume actually occupied by a chain, N/ρ 0 , and thus measures the degree of interdigitation of chains. At C → ∞, the SCF approximation becomes exact. The numerical simulations shown in Figure 3.6 were carried out at C = 50 (using the length of the copolymers as the reference length), which still seems large. Nevertheless, the effect of fluctuations is already quite dramatic. √ In three dimensions (d = 3), the Ginzburg parameter is proportional to the square root of the chain length, N . This is why the mean-field theory becomes very good for systems of polymers √ with high molecular weight, and only fails at selected points in the phase diagram. The relation C ∝ N has motivated the definition of an ‘invariant polymerization index’ N¯ = N b6 ρ02 ∝ C 2 , which is also often used to quantify fluctuation effects [72, 91]. In two dimensional systems, C is independent of the chain length and fluctuation effects are much stronger. Furthermore, topological constraints become important (i.e., the fact that chains cannot cross each other) which are not included in the Helfand model. In three dimensional systems of linear polymers, they only affect the dynamics (leading to reptation), but in two dimensions, they also change the static properties qualitatively [15]. Finally, we note that fluctuations can also be treated to some extent within the SCF theory, by looking at Gaussian fluctuations about the SCF solution [117]. This is useful for calculating structure factors and carrying out stability analyses. However, Gaussian fluctuations alone cannot bring about the qualitative changes in the phase behavior and the critical exponents which have been described above.

3.3.3

Analytical Theories

The SCF equations have to be solved numerically, which can be quite challenging from a computational point of view. In addition, they also serve as a starting point for the derivation of simpler approximate theories, which may even have analytical solutions in certain limits.

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Two main regimes have to be distinguished here. In the weak segregation limit, χ N is small, and the A and B homopolymers or copolymer blocks are barely demixed. This is the realm of the ‘random phase approximation’ (RPA), which can be derived systematically from the SCF theory. In the strong segregation limit, χ N is large, the polymers or copolymer blocks are strongly demixed and the system can basically be characterized in terms of its internal interfaces. 3.3.3.1

Weak Segregation and Random Phase Approximation

We first consider the situation at low χ N. In this case, the composition varies smoothly, A-rich domains still contain sizeable fractions of B-monomers and vice versa, and the interfaces between domains are broad, i.e., their width is comparable to the radius of gyration of the chains. The idea of the RPA is to perform a systematic expansion about a homogeneous reference state. More precisely, we use the SCF free density functional, Eq. (3.25), as a starting point, and then expand F about the homogeneous state. Defining A = ρ A /ρ 0 as usual, using A + B ≡ 1, and introducing the Fourier representation (k) = dr eikr (r), we obtain a functional of the form ⎧ ρ0 ⎨ 1  FRPA [ A (r)] = Vρ0 f homo + k B T | A (k)|2 2 (k) N ⎩ 2V k =0 ⎫ ⎬ 1     × (k) (k ) (−k − k )  (k, k ) + · · · (3.33) A A A 3 ⎭ 6V 2 k,k =0 where f homo is the SCF free energy per chain in the reference system, and the coefficient  n depend on the direct monomer interactions and on the intrachain correlations of free ideal Gaussian chains. We focus on the leading coefficient. To calculate  2 for a given blend, we define the pair correlators K i j [71, 117] K i j (k) = ρˆi (k)ρˆ j (k)

1 ρ0 N

(3.34)

which give the density-density correlations in an identical blend of noninteracting, ideal Gaussian chains and can thus be expressed in terms of Debye functions gD (x) (Eq. (3.7)). For example, for binary blends of ¯ j ( j = A, B), the pair homopolymers with chain length Nj , gyration radii Rg, j , and mean volume fractions correlators are given by



N ¯ N ¯ 2 2 KBB = K AB = K B A = 0 (3.35) A g D k 2 Rg,A B g D k 2 Rg,B K AA = NA NB For pure diblock copolymer blends with A-fraction f , one gets



K B B = (1 − f )2 g D (1 − f ) k 2 Rg2 K A A = f 2 g D f k 2 Rg2



 1  2 2

g D k Rg − f 2 g D f k 2 Rg2 − (1 − f )2 g D (1 − f )k 2 Rg2 K AB = 2

(3.36)

Having calculated K i j , one can evaluate  2 according to [71, 117]  2 =

K A A + K B B + K AB + K B A − 2χ N K A A K B B − K AB K B A

 (3.37)

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The function  2 is particularly interesting, because it is directly related to the structure factor of the homogeneous phase, S(k) ∝  2 (k)−1. Hence the RPA provides expressions for structure factors which can be compared to small angle scattering experiments, e.g., to determine effective interaction parameters. This is probably its most important application. We will now discuss specifically the application of the RPA to binary homopolymer blends and to diblock copolymer blends. (i) Binary Homopolymer Blends and Flory-Huggins-de Gennes Functional According to Eqs. (3.35) and (3.37), the RPA coefficient  2 for binary blends is given by   N /N A N /N B

+

− 2χ N 2 (k) = (3.38) 2 2 ¯ A g D k 2 Rg,A ¯ B g D k 2 Rg,B We assume that the composition varies only slowly in the system (on length scales not much shorter than 2 2 Rg ), and expand  2 for small wave vectors. Using gD (x) ≈ 1 − x/3 and Rg, j = b j N j /6, and inserting our result in the RPA expansion (3.33), we obtain the free energy functional ⎧  ⎨ 1  1 1 FRPA [ A ] ≈ Vρ0 f homo + k B Tρ0 + − 2χ | A (k)|2 ¯ A NA ¯ B NB ⎩ 2V k =0 ⎫   ⎬ 1  1 b2A b2B | A (k)|2 k 2 + + (3.39) ¯A ¯B ⎭ 2V k =0 18  The first two terms in (3.39) correspond to the second order expansion of the integral ρ 0 dr f SCF ( A ), where f SCF ( A ) is the SCF free energy per chain in a homogeneous system with A-volume fraction A . It thus seems reasonable to replace them by the full integral. The last term is a square gradient term in real space. Together, one recovers the Flory-Huggins-de Gennes functional of Section 3.3.1.2, Eq. (3.13). (ii) Copolymer Melts, Leibler Theory, and Ohta-Kawasaki Functional In diblock copolymer blends, Eqs. (3.36) and (3.37) yield the RPA coefficient   G(1) − 2χ N (3.40) 2 (k) = G( f )G(1 − f ) − (G(1) − G( f ) − G(1 − f ))2 /4 with the short hand notation G( f ) = 2g D ( f k 2 Rg2 ). At low χ N,  2 (k) is positive. Upon increasing χ N, one encounters a spinodal line where  2 (k) becomes zero for some nonzero k = q* , and the disordered state becomes unstable with respect to an ordered microphase separated state. Since the function  2 (k) is spherically symmetric in k, it does not favor a specific type of order. The information on possible ordered states is contained in the higher order coefficients  n , most notably, in the structure of the cubic term,  3 . In a seminal paper of 1979, Leibler has carried out a fourth order RPA expansion and deduced a phase diagram which already included the three copolymer phases L, C, and S (Figure 3.3) [118]. Milner and Olmsted later showed that the Leibler theory is also capable of reproducing the gyroid phase [119]. The RPA phase diagram roughly coincides with the full SCF phase diagram, as established 1996 by Matsen and Bates [95], up to χ N < 12. Unfortunately, fluctuations have a massive effect on the phase diagram at these small χ N (see Figure 3.5), therefore the predictive power of the Leibler theory must be questioned. Nevertheless, it is useful for identifying potential ordered phases and phase transitions in

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copolymer systems. Generalized Leibler theories still prove to be efficient tools to analyze phase transitions in complex copolymer blends by analytical considerations [120]. Next we attempt to construct a simplified free energy functional for diblock copolymer melts, in the spirit of the Flory-Huggins-de Gennes functional. To this end, we again expand  2 in powers of k, as in (i). Compared to homopolymer blends, however, there is an important difference:  2 (k) has a singularity at k → 0 and diverges according to 2 (k) ≈

2

1 3 2 2 − f ) k Rg2

at

f 2 (1

k→0

(3.41)

The singularity accounts for the fact that large-scale composition fluctuations are not possible in copolymer blends, since the A- and B-blocks are permanently linked to each other. It ensures that the structure factor S(k), vanishes at k → 0, suppresses macrophase separation and is thus ultimately responsible for the onset of microphase separation in the RPA theory. A 1/k2 term like (3.41) in a density functional corresponds to a long-range Coulomb type interaction. This observation motivated Ohta and Kawasaki [121] in 1986 to propose a free energy functional for copolymer melts, which combines a regular squaregradient functional accounting for direct short-range interactions with a long-range Coulomb term accounting for the connectivity of the copolymers. In real space, the Ohta-Kawasaki functional has the form    ρ0 B ρ0 A 2 (3.42) dr W( A ) + (∇ A ) + dr dr G(r, r ) δ A (r) δ A (r ) FOK [ A ] = N 2 N 2 ¯ A . The last term introduces the long-range interactions, with G(r, r ) defined such with δ A (r) = A (r) − that G(r, r ) = −δ(r − r )

(3.43)

which corresponds to G(r, r ) ∼ 1/|r − r | in infinitely extended systems. Given Eq. (3.41), it seems natural to identify A = 3/(2 f 2 (1 − f )2 Rg2 ). The choice of W and B is somewhat more arbitrary. The function W( A ) is a free energy density with two degenerate minima and can be approximated by a fourth order polynomial in A . As for the coefficient of the square gradient term, B, Ohta and Kawasaki originally estimated it from the asymptotic behavior of  2 at k → ∞, 2 (k) ≈

1 k 2 Rg2 2 f (1 − f )

at

k→∞

(3.44)

which yields B = Rg2 /(2 f (1 − f )). Later, they noted that this choice of B gives the wrong interfacial width at stronger segregation, which has implications for the elastic constants and the equilibrium period of the ordered phases, and suggested to replace B by a constant in the strong segregation limit [122]. The Ohta-Kawasaki functional reproduces microphase separation and complex copolymer phases such as the gyroid phase [123, 124] and even the Fddd phase [124]. It can be handled much more easily than the Leibler theory or the full SCF theory (see Section 3.4.3.2), therefore it is particularly popular in largescale dynamical simulations of copolymer melts (see Section 3.4.3.2). Different authors have generalized it to ternary blends containing copolymers [123, 125, 126]. In particular, Uneyama and Doi have recently proposed a general density functional for polymer/copolymer blends that reduces to the Flory-Huggins-de Gennes functional in the homopolymer case and to the Ohta-Kawasaki functional in the diblock case.

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3.3.3.2

53

Strong Segregation

We turn to discussing the situation at high χ N. The A-rich and B-rich (micro)phases are then well-separated by sharp interfaces. The free energy contribution from the interfacial regions (i) and the chain conformations inside the A- or B-domains (ii) can be treated separately. (i) Interfacial Profiles and Ground State Dominance In the interfacial region, the free energy is dominated by the contribution of the direct A-B interactions and the local stretching of segments. Chain end effects can be neglected. This simplifies the situation considerably. We first note that the diffusion equation ((3.28) or (3.29)) for a j-chain or a j-block of a chain has the same structure than the time-dependent Schr¨odinger equation, if one identifies s ↔ it. As is well known from quantum mechanics, the general solution can formally be expressed as [127] q j (r, s) = n cn ψn, j (r)e−n, js , where {ψ n,j (r),  n,j } are Eigenfunctions and Eigenvalues of the operator (b2j /6  − W j (r)/N ). At large s, the smallest Eigenvalue  0,j dominates, i.e., q j (r, s) ∝ ψ0, j (r)e−0, j s , and the resulting density in the large Nj limit is ρ j ∝ |ψ0, j |2 e−0, j N j . This type of approximation is called ‘ground state dominance’. It is commonly used to study polymers at interfaces and surfaces. In the case of blends, we have the freedom to shift the fields Wj (r) by a constant value, hence we can set  0,j = 0. The self-consistent field equations can thus be written as  ρ j = ρ0 |ψ j (r)|

2

with

b2j

Wj − 6 N

 ψj = 0

and

Wj =

N δH I −ξ k B T δρ j

(3.45)

 where ψ j is normalized such that Q j = dr|ψ j (r|2 = V j is the partial volume occupied by the polymers j,  and ξ (r) ensures j |ψ j |2 ≡ 1. In order to derive an epression for the free energy, we first note that Eqs. (3.45) minimize a Lagrange action, 

kB T ρ0 L = HI + (3.46) dr b2A (∇ψ A )2 + b2B (∇ψ B )2 6 with respect to ψ j under the constraint |ψ A |2 + |ψ B |2 ≡ 1. One easily checks that L vanishes for homogeneous bulk states, and that the minimized L is equal to the extremized SCF Hamiltonian FSCF , Eq. (3.25) up to a constant. Hence L can be identified with the interfacial free energy. Rewriting it in terms of the volume fractions φ j and using (∇ A )2 = (∇ B )2 , one obtains the free energy functional  Fint [ A (r)] = H I + ρ0

dr

kB T 24



b2A b2 + B A B

 (∇ A )2

(3.47)

which reproduces Eq. (3.14) for Flory-Huggins interactions (3.17). For bA = bB = b and Flory-Huggins interactions, the self-consistent field equations √ (3.45) are solved by a tanh profile, (ρ A − ρ B√) ∼ ρ0 tanh(z/wSSL ) with the interfacial width wSSL = b/ 6χ and the interfacial tension σSSL = k B Tρ0 b χ /6 [66–68]. (ii) Copolymer Conformations and Strong Stretching Theory The free energy functional (3.47) is sufficient to describe strongly segregated homopolymer blends. In copolymer blends, additional contributions come into play due to the fact that the copolymer junctions are confined to the interfaces and the copolymer blocks stretch away from them into their respective A or B domains. The associated costs of configurational

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free energy can be estimated within a second approximation scheme, the ‘strong stretching’ theory (SST) [128–131]. The main idea of the SST was put forward in 1985 by Semenov [128], who noted that for strongly stretched copolymer blocks, the paths fluctuate around a set of ‘most probable paths’. This motivates to approximate the single-chain partition function Q, Eq. (3.21), by its saddle point, i.e., the path integral in Q is replaced by an integral over ‘classical’ paths Rc that extremize the integrand and thus satisfy the differential equation [132] 3 d2 Rc 1 = ∇W (Rc ) 2 2 b ds N

(3.48)

We will treat the copolymer blocks as independent chains of length M. The classical paths corresponding to one block are then characterized by their boundary conditions, Rc (0) = r j and Rc (M) = re , where the junction rj is confined to an interface and the free end re is distributed everywhere in its domain. Next, we note that for infinitely long blocks, the classical paths must satisfy dRc /ds|s=M = 0

M →∞

for

(3.49)

at the free end. Mathematically speaking, they would not have a well-defined end position otherwise. Physically speaking, the ‘average’ chain representing the classical path does not sustain tension at the free end, which seems reasonable. In the following, Eq. (3.49) is also imposed for finite (large) blocks as an additional boundary condition. Eq. (3.48) is then overdetermined and can, in general, no longer be solved for arbitrary end positions re . To ensure that chain ends are indeed free to move throughout the domain, the field W(r) must have a special shape. Specifically, near flat interfaces it must be parabolic as a function of the distance z from the interface [129, 130], 3 π2 2 1 W (z) = − z N 8 b2 M 2

(3.50)

This is one of the main results of the SST. It generally applies to situations where strongly stretched polymers are attached to an interface, e.g., strongly segregated copolymer blocks, [72, 133] or polymer brushes in solvents of arbitrary quality [131]. The SST field must always have the form (3.50), and the remaining task is to realize this by a suitable choice of the chain end distribution P(re ). In the incompressible blend case, P(re ) must be chosen such that the density in the domains is constant, ρ 0 . Luckily, we do not have to evaluate P(re ) explicitly to calculate the free energy. The SST field has another convenient property: One can show that the stretching energy of classical paths of fixed length N in a field satisfying Eqs. (3.48) and (3.49) is exactly equal to the negative field energy, 

N 0

3 ds 2 2b



dRc ds

2

1 =− N



N

ds W (Rc (s))

(3.51)

0

Summing over all blocks in a domain, the total stretching energy is thus given by 1 Fstretch =− kB T N

 dr ρ(r) W (r) ≈ ρ0

3 π2 8 b2 M 2

 dr z 2

(3.52)

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55

where the integral is over the volume of the domain, and z denotes the closest distance to an interface. The total energy of the system can be estimated as the sum over the stretching energies in the different domains, Eq. (3.52), and the interfacial energy, Eq. (3.47), and then used to evaluate the relative stability of different phases. In the strong segregation limit, only the C, L, and S phase are found to be stable [72], in agreement with SCF calculations at high χ N. The validity of the strong stretching theory seems to be restricted to very large chains [134]. This is presumably to a large extent due to the requirement (3.49), which does not necessarily hold for classical paths of finite length. Netz and Schick [135, 136] have shown that an unrestricted ‘classical theory’, which just builds on the saddle point integration of Q and avoids using (3.49), gives results that agree better with the SCF theory. However, the classical theory has to be solved numerically, and the computational advantage over the full SCF theory is not evident. The SST has found numerous applications [72] and has been extended and improved in various respect. It provides an analytical approach to analyzing multicomponent polymer blends in a segregation regime where the SCF theory becomes increasingly cumbersome, due to the necessity of handling narrow interfaces. 3.3.4

An Application: Interfaces in Binary Blends

To close the theory section, we discuss the simplest possible examples of an inhomogeneous polymer system: An interface in a symmetrical binary homopolymer blend. This system has been studied intensely in experiments [137–143]. By mixing random copolymers of ethylene and ethyl-ethylene with two different, but very well defined copolymer ratios, Carelli et al. [137, 143] were able to tune the Flory-Huggins parameter very finely and study interfacial properties in a wide range of χ N between the weak segregation limit and the strong segregation limit. Figure 3.7 compares the results for the interfacial width and compares them with the mean-field prediction for the weak segregation limit, the strong segregation limit, and the full numerical result.

500

Interfacial width 2w (Å)

400

300

200

100

0 2

8

14

20

26

32

38

χN

Figure 3.7 Intrinsic width of interfaces between A- and B-phases in binary polyolefin blends as a function of χ N, compared with the predictions of the weak segregation theory, the strong segregation theory, and the full SCF theory. From Ref. [137].

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We should note that there is a complication here. As we have mentioned earlier, fluid–fluid interfaces are never flat, they exhibit capillary waves [144, 145]. This leads to an apparent broadening of the interfacial width w [146]. The apparent width depends on the lateral length scale of the observation L and can be calculated according to [147, 148] w = 2

w02

  L kB T ln + 4σ a0

or

w = 2

w02

  L 2k B T ln + πσ a0

(3.53)

depending how it is measured. Here w0 is the ‘intrinsic’ width, σ the interfacial tension, and a0 a ‘coarsegraining length’, which is roughly given by the interfacial width [146, 148]. Both the quantities L and a0 are not very well defined in an actual experiment. Fortunately, they only enter logarithmically, therefore the result is not very sensitive to their values. The theoretical curves shown in Figure 3.7 include the capillary wave broadening, calculated using the interfacial tension from the respective theory. Comparing the curves in Figure 3.7, one finds that the weak segregation theory consistently overestimates the width, and the strong segregation theory consistently underestimates it. The numerical SCF values interpolate between the two regimes and are in excellent quantitative agreement with the experimental data over almost the whole range of χ N. The SCF theory is also found to perform very well compared to computer simulations [148, 149]. It reproduces many features of the interfacial structure, such as chain end distributions, local segment orientations, etc. at a quantitative level, if capillary waves are accounted for [148]. This illustrates the power of the SCF theory to describe the local structure of inhomogeneous polymer systems, even if the global structure is affected by large-scale composition fluctuations.

3.4

Simulations of Multiphase Polymer Systems

Whereas theoretical work on multiphase polymer systems has a long-standing tradition, the field of simulations in this area is much younger. This is because polymer simulations are computationally very expensive, which has essentially rendered them unfeasible until roughly 20 years ago. In this section, we will attempt to give an overview over the current state-of-the-art of simulations of inhomogeneous multicomponent polymer systems.

3.4.1

Coarse-Graining, Part II

One of the obvious challenges in multiphase polymer simulations is that polymers are such big molecules, which moreover self-organize into even larger supramolecular structures. Polymeric materials exhibit structure on a wide range of length scales, from the atomic scale up to micrometers. Their specific material properties are to a great extent determined by local inhomogeneities and internal interfaces, and depend strongly on the interplay between these mesostructures in space and time. In order to understand the materials and make useful predictions for new substances, one must analyze their properties on all time and length scales of interest. Therefore, multiscale modeling has become one of the big topics in computational polymer science. The central element of multiscale modeling is coarse-graining. By successively eliminating degrees of freedom (electronic structure, atomic structure, molecular structure, etc.), a hierarchy of models is constructed (see Section 3.2.2). For each type of model, optimized simulation methods are developed, which allow one to investigate specific aspects of the materials. Having identified suitable classes of coarse-grained models, one can proceed in two different manners.

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(i) Generic Modeling This approach has been favored historically, and up to date, the overwhelming majority of simulations of multicomponent polymer systems is still based on it (see Section 3.4.4). Generic models are simple and computationally efficient. They are not designed to represent specific materials; rather, they are the simulation counterpart of the theoretical models discussed in the previous section. They are suited to study generic properties of polymer and copolymer systems, i.e., to identify the behavior that can be expected from their stringlike structure, their chemical incompatibility etc. Simulations of generic models are also particularly suited to test theories. Generic models are used in all areas of materials science, and in most cases, they only give qualitative insights into the behavior of a material. This is different for polymers, because of their universal properties (see Section 3.2). For example, we have already seen in Figure 3.7 that a generic theoretical model (the Edwards model) quantitatively predicts important aspects of the interfacial structure in real polymer melts. Nevertheless, the predictive power of generic models is restricted, and relies on the knowledge of ‘heuristic’ parameters such as the χ -parameter. Therefore, a second approach is attracting growing interest.

(ii) Systematic Bottom-Up Modeling The idea of systematic coarse-graining is to establish a hierarchy of models for the same specific material, starting from an ab-initio description, with well-defined quantitative links between the different levels. Ideally, the goal is to replace many degrees of freedom by a selection of fewer ‘effective’ degrees of freedom. If one is only interested in equilibrium properties, the problem is at least well defined. For each possible coarse-grained configuration, one must evaluate a partial partition function of the full system under the constraints imposed by the values of the coarse-grained degrees of freedom. This procedure results in an effective potential in the coarse-grained space, which is, in general, a true multibody potential – it cannot be separated into contributions of pair potentials. If one is interested in dynamical properties, the situation is even more complicated. One must replace a dynamical system for all variables by a lower dimensional system for a subset of effective variables. This can be done approximately, e.g., using Mori Zwanzig projector operator techniques [150]. The new dynamical system is inevitably a stochastic process with memory, i.e., the future time evolution not only depends on the current state of the system, but also on its entire history. Obviously, such ‘ideal coarse-graining’ is not feasible for polymer systems. Instead, researchers adopt a heuristic approach, where they first define a coarse-grained model, which typically has no memory and only pair potentials, and match the properties of the coarse-grained model with those of the fine-grain model as best they can: The model parameters are chosen such that the coarse-grained model reproduces physical properties of interest, such as correlation functions or diffusion constants [151–154]. Already at an early stage, researchers have started to develop schemes for mapping real polymers on lattice models [155–157]. Nowadays, off-lattice models are more common. Early approaches focussed on the task of reproducing the correct intrachain correlations by optimizing the bond potentials in the chains [155]. Later, the interchain correlations were considered as well, which can be matched by adjusting the non-bonded, intermolecular potentials in the coarse-grained model. It is important to note that the resulting effective potentials depend on the concentration and the temperature [158] (much like the χ -parameter itself). Different methods to determine effective potentials have been devised and even automated packages are available [153, 159–162]. The reverse problem – how to reconstruct a fine-scale model from a given coarsescale configuration – has also been addressed [163, 164]. Nowadays, the available techniques for mapping static properties are relatively advanced. In contrast, the field of mapping dynamical properties is still in its infancy [165]. The standard multiscale approach is sequential, i.e., numerical simulations are carried out separately for different levels. Currently, increasing effort is devoted to developing hybrid schemes where several coarsegraining levels are considered simultaneously within one single simulation [166, 167].

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Despite the large amount of work that has already been devoted to systematic coarse-graining, coarsegrained simulation studies of realistic multicomponent polymer blends are still scarce. Mattice and coworkers have carried out lattice simulations of blends containing polyethylene and polypropylene homopolymers and copolymers [168–170]. Faller et al. have developed and studied a coarse-grained model for blends of polyisoprene and polystyrene [171–173]. 3.4.2

Overview of Structural Models

After these general remarks, we shall give a brief overview over the different models that are currently used in multicomponent polymer simulations. 3.4.2.1

Atomistic Models

Atomistic simulations are computationally intensive, and rely very much on the quality of the force fields. (Force fields are a separate issue in multiscale modeling, which shall not be discussed here.) Therefore, atomistic simulations of blends are still relatively scarce. So far, most studies have focussed on miscibility aspects [174–184]. Already early on, atomistic and mesoscopic simulations were combined in multiscale studies: Atomistic simulations were used to determine the Flory-Huggins χ -parameter, coarse-grained methods were then applied to study large-scale aspects of phase separation [185–191] or mesophase formation [192]. Only few fully atomistic studies deal with aspects beyond miscibility, e.g., the formation of lamellar structures in diblock copolymers [193], or the diffusion of small molecules in blends [194]. 3.4.2.2

Coarse-Grained Particle Models

The coarse-grained models for polymers can be divided into two main classes: Coarse-grained particle models operate with descriptions of the polymers that are considerably simplified, compared to atomistic models, but still treat them as explicit individual objects. Field models describe polymer systems in terms of spatially varying continuous fields. We begin with discussing some of the most common particle models. Lattice Chain Models Lattice models have the oldest tradition among the coarse-grained particle models for polymer simulations, and are still very popular. The first molecular simulations of multiphase polymer systems – studies of binary homopolymer blends by Sariban and Binder in 1987 [195, 196] and by Cifra and coworkers in 1988 [197] – were based on lattice models. They are particularly suited to be studied with Monte Carlo methods, and several smart Monte Carlo algorithms have been designed especially for lattice polymer simulations [198, 199]. In molecular lattice models, the polymers are represented as strings of monomers confined to a lattice. A natural approach consists in placing the ‘monomers’ on lattice sites and linking them by bonds that connect nearest-neighbor sites. For many applications, it has proven useful to apply less rigid constraints on the links and allow for bonds of variable length, which may also connect second-nearest neighbors [200] or stretch over even longer distances [201]. Moreover, the lattice is usually not entirely filled with monomers, but also contains a small fraction of voids. This is because most Monte Carlo algorithms for polymers do not work at full filling, and special algorithms have to be devised for that case [202]. One particularly popular lattice model is the ‘bondfluctuation model’, devised in 1988 by Carmesin and Kremer [203]. It is based on the cubic lattice; monomers do not occupy single sites, but√entire cubes in a cubic lattice. They are connected by √ √ ‘fluctuating bonds’ of varying length, (2, 5, 6, 3 or 10 lattice constants). In the bond-fluctuation model, a polymer system behaves like a dense polymer melt already at the volume fraction 0.5. Therefore, it can be simulated very efficiently.

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Despite their intrinsically anisotropic character, lattice models are able to reproduce most known selfassembled mesophases in copolymer melts, even the gyroid phase in diblock copolymers [204]. Nowadays, they are used to study such complex systems as ABC triblock copolymer melts confined in cylindrical nanotubes [205], which feature a rich spectrum of novel morphologies, e.g., stacked disks, curved lamellar structures, and various types of helices. Off-Lattice Chain Models For many years, only lattice simulations were sufficiently efficient that they could be used to study polymer blends at a molecular level. With computers becoming more and more powerful, off-lattice chain models have become increasingly popular. Compared to lattice models, they have the advantage that they provide easy access to forces and can also be used in Molecular Dynamics or Brownian Dynamics simulations. They do not impose restrictions on the size and shape of the simulation box (in lattice models, the box dimensions have to be multiple integers of the lattice constant). The structure of space is not anisotropic as in lattice models. Whereas the inherent anisotropy of lattice models does not seem to cause problems if the lattice model is sufficiently flexible and if the chains are sufficiently long, simulations of shorter chains can be hampered by lattice artifacts. In bead-spring models, polymers are represented by chains of spherically symmetric force centers connected by springs. Numerous variants have been proposed, which differ in the choice of the spring potentials (the bonded interactions) and the choice of the pairwise interactions between the beads (the non-bonded potentials). The simplest choice of spring potential is a harmonic potential. In order to prevent chains from crossing each other in dynamical simulations, an anharmonic cutoff on the spring length is often imposed. A popular choice is the ‘Finitely Extensible Nonlinear Elastic’ (FENE) potential   k 2 (b − b0 )2 VFENE (b) = d ln 1 − 2 d2

(3.54)

which reduces to a harmonic spring potential with equilibrium spring length b0 at b ≈ b0 , and diverges at |b − b0 | → d. In some applications, the springs are constrained to have fixed lengths – however, this requires the use of special algorithms and changes slightly the distribution of bond angles [198]. In addition, some bead-spring models also include bending potentials that allow one to tune the chain stiffness, or even torsional potentials. The non-bonded interactions drive the segregation of the monomers. As we have discussed in Section 3.3.1.3, both energetic and entropic factors can contribute to making monomers incompatible. Many models operate with energetic monomer (in)compatibilities, but models with entropically driven segregation are also common. As an example, we consider one commonly used type of potential, the truncated Lennard-Jones potentials   for r < rc VLJ (r ) =  (σ/r )12 − (σ/r )6 + C

(3.55)

(V LJ = 0 otherwise), where the parameter C is chosen such that V LJ (r) is continuous at r = rc . If the cutoff parameter rc is larger than 21/ 6 σ (a common choice is r = 2.5σ ), the potential has a repulsive core surrounded by an attractive well. In this case, energetic incompatibility can be imposed by using species dependent interaction parameters  ij with 2 AB <  AA +  BB . If the cutoff parameter is rc ≤ 21/6 , the potential is purely repulsive. In this case, one can still enforce monomer segregation by choosing species dependent and non-additive interaction radii σ ij with 2σ AB >σ AA + σ BB . The mechanism driving the segregation is the Equation-of-State effect discussed in Section 3.3.1.3. A simulation model that is based on this idea has been

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proposed by Grest, Lacasse and Kremer in 1996 [206]. It is more efficient than conventional models with energetically driven segregation, because the interaction range is shorter. We note that high pressures have to be applied to keep the chains together and to drive the demixing. In the context of ‘Dissipative particle dynamics’ (DPD) simulations, (see below) it has also become popular to use soft non-bonded potentials without a hard core. A typical DPD-potential has the form [207] VDPD (r ) =

v (1 − r/rc )2 2

for r < rc

(3.56)

(V DPD = 0 otherwise). Demixing is induced by using species dependent parameters vi j > 0 with 2v AB > (v A A + v B B ). The mechanism driving the segregation is again related to the Equation-of-State effect – like particles overlap more strongly than unlike particles. The DPD simulation method was originally proposed in the context of fluid simulations, where every DPD particle supposedly represents a lump of true particles. This motivates the absence of hard, impenetrable cores and even a roughly linear shape [208]. Early on, DPD potentials were also used in polymer simulations [207]. As long as topological constraints are not important, the monomer potentials do not need to have a hard core (see also the discussion at the end of the next section). ‘Edwards’ Models A special class of chain models are the Edwards models, which shall be treated separately. The idea is to implement directly an Edwards-type interaction H I (Eq. (3.9)) in a particle-based simulation. The molecules are modeled as off-lattice chains as before, and the non-bonded interactions are given by a potential V = H I [{ρi }] that depends on local monomer densities ρi (r). To complete the definition of the model, one must prescribe how to evaluate the local densities. This is most easily done by simply counting the monomers in each cell of a grid. Other prescriptions that do not impose a grid are also conceivable. Edwards models are not yet very common in simulations of multiphase polymer systems, but they will very likely gain importance in the future. The basic idea was put forward by Zuckermann and coworkers [209, 210] in 1994 in the context of polymer brush simulations. It was first applied to studying microphase separation in block copolymers by Besold et al. [211], who also showed that the model produces correct single-chain behavior in solution [212] (i.e., the chains behave like regular self-avoiding chains). In the following, the power of the approach has been demonstrated in a series of impressive work by M¨uller and coworkers [213] (see also Section 3.4.3.1). Two notes are in order here. First, it is difficult to impose strict incompressibility in particle-based simulations. Instead, dense melt simulations usually operate at finite compressibility: One introduces an additional term [66] H I,comp. =

κ (ρ A + ρ B − ρ0 )2 2

(3.57)

with a high modulus κ in the interaction Hamiltonian. Second, chains may overlap in the Edwards models. They share this ‘problem’ with the DPD models introduced in the previous section. In three dimensions, the lack of hard core interactions has no effect on the static properties of linear polymers [212, 213]. Topological constraints become important in two dimensions [15], or in melts of cyclic polymers [214], or, most notably, when looking at dynamical properties. In Edwards models, chains do not entangle, and reptation dynamics has to be put in ‘by hand’ [215]. Ellipsoid Model In 1998, Murat and Kremer proposed a model that allows to study weakly segregated polymer blends on the scale of the gyration radius Rg and beyond [216]. If details of the conformations are not of interest, the polymers can be replaced by single, soft particles with ellipsoidal shape. The idea was

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then pursued mainly by Eurich, Maass, and coworkers [217–219], who also suggested to extend the model to diblock copolymers by modeling them as dimers to account for their dumbbell shape [219]. Sambrisky and Guenza have recently worked out a microscopic foundation for coarse-graining diblock copolymers into such dumbbells, which is based on liquid-state integral equations [220]. 3.4.2.3

Field-Based Models

Field-based models for polymer systems no longer treat the molecules as individual objects, but describe them in terms of locally varying fluctuating fields. Among these, the molecular field theories still incorporate some information on the conformations of molecules, and the Ginzburg-Landau models focus on the large-scale structure only. Field-Theoretical Models Field-theoretical models are the molecular field equivalent of the particle-based ‘Edwards’ models discussed above. The idea is to use the starting point of the selfconsistent field theory, i.e., a field-theoretic expression for the partition function (e.g., Eq. (3.19)), and to evaluate the functional integrals over fluctuating fields by simulation methods. Since some fields are complex, one is confronted with a sign problem (the integrand oscillates between positive and negative). Fredrickson and coworkers have demonstrated that the integrals can nevertheless be evaluated in many cases using a method borrowed from elementary particle physics, the ‘Complex Langevin’ simulation method [8, 221, 222]. We refer to Ref. [8] for a general presentation of the method, and Ref. [223] for technical details. To be more specific, let us consider the system discussed in Section 3.3.2.1, a blend of polymers/copolymers containing two types of monomers A and B, with Flory-Huggins interactions (3.17). Rather than evaluating the integral over all five fluctuating fields WA,B , ξ , ρ A,B in Eq. (3.19), we reconsider the original partition function, Eq. (3.18), and decouple the integrals over different paths (polymer conformations) by means of a Hubbard-Stratonovich transformation [90]. This is possible, because the Flory-Huggins interaction H I [{ρˆi }] is quadratic in the densities ρˆi . The result is a functional integral over ‘only’ two fields – one which is conjugate to the total density ρˆ = ρˆ A + ρˆ B and imaginary, and one which is conjugate to the composition mˆ = ρˆ B − ρˆ A and real: 

 DW+

Z∝ i∞

∞

DW− e−FFTS /k B T

(3.58)

with ρ0 FFTS [W+ , W− ] = N



1 χN



 dr W−2 −

dr W+ −

 α

N Vα ln ρ0 Qα /n α Nα

 (3.59)

This partition function is the starting incompressible AB (co)polymer systems. The idea to study multiphase polymer systems by direct evaluation of fluctuating field integrals like Eq. (3.58) was first put forward by Ganesan and Fredrickson in 2001 [224]. They used Complex-Langevin simulations to look at fluctuation effects in pure symmetric diblocks. Since then, Fredrickson and coworkers have shown that the method can be extended in various ways, e.g., it can deal with external stresses [225]), and it can be applied very naturally to charged polymers [226, 227]. Nevertheless, there are still problems with it. The theoretical foundations of the Complex Langevin method are not fully established. One can show that under certain conditions, it produces the correct statistical averages if the system reaches equilibrium [8]. However, the stability of the simulations cannot be ensured. Even if the Langevin step size chosen is small, a

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fraction of simulation runs still crashes [223]. Such problems are unknown from other simulation methods. Because it is generally not widely used and very young in the polymer community, the Complex Langevin method still suffers from teething troubles. An alternative approach to carrying out field-theoretical simulations in melts has been suggested by Do¨uchs et al. [110] (see also Ref. [91]). They suggested to treat only the real integral over W − with computer simulation and approximate the problematic imaginary integral over W + by its saddle-point. The integral over W − can be evaluated by standard Monte Carlo methods. The results from this partial saddle point approach were shown to agree quantitatively with the full Complex Langevin simulation [110]. Possible strategies to improve upon the saddle point integral have been pointed out by B¨aurle and coworkers [228–230]. As already pointed out, field-theoretic models are in some sense equivalent to particle-based Edwards models (if the latter use discrete Gaussian chain models). In field-theoretical simulations, the mean-field limit can be reached very naturally and at low computational cost. Hence they are more efficient than particle-based simulations at high polymer densities, close to the SCF limit, whereas particle-based simulations are superior at lower polymer densities. An application of field-theoretic simulations has already been shown in Section 3.3.2.5. The simulation data shown in Figure 3.6 were obtained with field-theoretic Monte Carlo simulations. Molecular Density Functionals Molecular density functionals are free energy functionals of the type (3.25). They are used in dynamical density functional simulations (see Section 3.4.3.2). Ginzburg-Landau Models Ginzburg-Landau models no longer incorporate specific information on chain conformations and thus have a much simpler structure than molecular field models. Examples are the FloryHuggins-de Gennes functional for homopolymer blends (Eq. (3.13)), or the Ohta-Kawasaki functional for diblock copolymer melts (Eq. (3.42)). This is the highest level of coarse-graining discussed in the present article. 3.4.3

Overview of Dynamical Models

In many cases, only static equilibrium properties are of interest, and then most dynamical models are equivalent. When looking at dynamical properties such as dynamical response functions, or at nonequilibrium situations, the choice of the dynamical model becomes relevant. In the following, we summarize some important models that have been used to study multiphase polymer systems. 3.4.3.1

Particle-Based Dynamics

Kinetic Monte Carlo (MC) A priori, the Monte Carlo (MC) method has been invented as a method to evaluate high dimensional integrals (i.e., thermal averages), and is designed for studying dynamics. Nevertheless, MC simulations are used for dynamical studies, based on the fact that like many static properties, dynamical phenomena are also often governed by universal principles. In kinetic MC simulations, one analyzes the artificial Monte Carlo evolution of a system in order to gain insight into real dynamical processes in the system. The main requirement is that the Monte Carlo moves are only local and reasonably ‘realistic’, i.e., chain crossings are not allowed. Kinetic Monte Carlo simulations have been used, e.g., to study the early stages of demixing in binary blends or the ordering dynamics in block copolymer melts. Molecular Dynamics (MD) In Molecular Dynamics simulations, one solves directly Newton’s or Hamilton’s equations of motions. This is the most straightforward approach to studying dynamical processes in a system. As usual, of course, the devil is in the detail [198].

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Brownian Dynamics In Brownian Dynamics simulations, one also solves equations of motion, but the underlying dynamics is not Hamiltonian. As we have noted earlier (Section 3.4.1), systematic coarse-graining necessarily turns a dynamical system into a stochastic process. Brownian Dynamics simulations account for this fact. They include dissipative and stochastic forces, which supposedly represent the interaction of the coarse-grained degrees of freedom with those that have been integrated out. In general, however, these forces are not derived systematically from the original model, but postulated heuristically. More specifically, the particles j experience three types of forces: m j r¨ j = FCj + F Dj + F Rj

(3.60)

The first term FCj correspondes to the conservative forces, which are derived from the interparticle potentials of the structural model under consideration. These forces also enter the standard Molecular Dynamics simulations. The second F Dj term is a dissipative force that couples to the velocity of the particles. The last term F Rj is a Gaussian stochastic force with mean zero whose amplitude is related to the dissipative force by a ‘fluctuation-dissipation theorem’ [231]. In a canonical simulation, the last two terms constitute a thermostat, i.e., they maintain the system at a given temperature T. Microcanonical models where energy is randomly shifted between ‘internal’ and ‘external’ degrees of freedom have been designed as well. In the simplest (canonical) Ansatz, the dissipative force on a particle j at time t is proportional to its velocity v j (t). According to the fluctuation-dissipation theorem, the stochastic force then fulfills F Dj = −γ j v j

←→



 R R F j,α (t) Fk,β (t  ) = 2γ j k B T δαβ δ jk δ(t − t  )

(3.61)

Other choices are possible. For example, the dissipative force on j can depend on the velocities of other particles k as well, and/or on the history of the system. The fluctuationdissipation relation for the stochastic force has to be adjusted accordingly [231]. The dynamics defined by Eq. (3.61) is commonly used because it is so simple, but it has the disadvantage that it is not Galilean-invariant – a system moving at constant speed is treated differently than a system at rest. Therefore, it does not incorporate hydrodynamic effects correctly, and it is not suited to study nonequilibrium systems such as polymer melts in shear flow. Dissipative Particle Dynamics (DPD) The problems of Eq. (3.61) are avoided in the recently developed ‘Dissipative Particle Dynamics’ (DPD) method [232, 233]. DPD is a special type of Brownian Dynamics which is Galilean-invariant and has become very popular in simulations of complex fluids in recent years. The dissipative and stochastic forces are constructed such that they conserve the total momentum and angular momentum in the system. Consequently, they couple to relative velocities of particles rather than absolute velocities, and they act as central forces (to preserve angular momentum). Specifically, the forces acting on a particle j are given by FCj −



γ ω(r jk ) (ˆr jk v jk ) rˆ jk +

!

2γ k B T ω(r jk ) rˆ jk ζ jk

" (3.62)

k = j

where r jk = r j − rk is the vector separating two particles j and k, rjk its length, rˆ jk the unit vector in the same direction, ω(r) an arbitrary function with a cutoff, and ζ jk are symmetric and uncorrelated Gaussian random numbers with mean zero and variance one. In the literature, the term ‘DPD simulations’ often refers to simulations that use DPD dynamics in combination with soft ‘DPD potentials’ (see Section 3.4.2.2, (Eq. (3.56)). However, DPD can of course be used as a dynamical model [233] or simply as a thermostat [234] in combination with any type of potential.

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In contrast to kinetic Monte Carlo simulations or simple Brownian dynamics simulations (using Eq. (3.61)), DPD simulations take full account of hydrodynamic interactions. Studies of microphase separation in copolymer melts have shown that this makes a difference. The dynamics is strongly affected by hydrodynamic effects in certain regions of phase space – in particular, hydrodynamic interactions play a critical part in helping the system to escape from metastable transient states [235]. Single Chain in Mean Field (SCMF) In 2005, M¨uller proposed an efficient method to study the ordering dynamics of polymer blends within Edwards models (see Section 3.4.2.2) [236, 237]. The idea is to take snapshots of the density configurations in regular intervals, and to let the chains move in the fields created by these snapshots, e.g., by kinetic Monte Carlo simulations. If the fields were updated after every Monte Carlo move, this would correspond to a regular simulation of the Edwards model. Daoulas et al. have shown that it is possible to update much less often [237] without changing the results. This makes the method very efficient and especially suited for the use on parallel computers. A similar idea had been put forward already in 2003 by Ganesan and Pryamitsin [238], in a less transparent formulation that involves self-consistent fields, to study stationary inhomogeneous polymer systems in an externally imposed flow field. In the absence of flow, the model of Ganesan and Pryamitsin is equivalent to that of M¨uller et al. A hybrid method that combines kinetic Monte Carlo and the self-consistent field formalism has also recently been proposed by B¨aurle and Usami [239]. SCMF simulations are hampered by the general limitations of the Edwards models – chains can cross each other. However, there are ways to introduce the dynamical effect of entanglements at least in cases where the equilibrium configuration space isnot affected by topological constraints [215]. 3.4.3.2

Field-Based Dynamics

Dynamic Density Functional Theory (DDFT) Dynamic density functional theories (DDFT) are based on density functionals for polymer systems, such as, e.g., Eq. (3.25). They supplement them by a model for their dynamical evolution at nonequilibrium. For diffusive dynamics, the dynamical equations in an imposed flow profile v have the general form. [105–108] ∂ρi + ∇(vρi ) = ∂t



dr

 ij

∇r i j (r, r )∇r

δF + ηi (r, t) δρ j (r )

(3.63)

where i j (r, r ) is a generalized mobility, and ηi (r, t) a Gaussian white noise with mean zero. If the amplitude of the latter is very small or zero, one has ‘mean-field dynamics’ and the system evolves towards a minimum of the free energy functional (although not necessarily the global minimum). If the noise is larger and satisfies the fluctuation dissipation theorem, ηi (r, t)η j (r , t  ) = −2k B T δ(t − t  ) ∇r i j (r, r )∇r

(3.64)

the density configurations {ρ i (r)} in an equilibrium simulation (no flow, sufficiently long runs) will be distributed according to P[ρ] ∝ exp(−F[ρ]/k B T ). The kinetic Onsager coefficient i j (r, r ) depends on the microscopic dynamics in the system. Since it characterizes the current of component i in response to an external force acting on component j, it is reasonable to assume that it is proportional to the local density ρ i (r). An efficient choice which however disregards the connectivity of the chains is thus i j (r, r ) = Mi ρi (r) δ(r − r )δij for compressible systems [240]. or (r, r ) = Mρ Aρ B δ(r − r ) for binary incompressible systems (local dynamics) [106].

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To account for the chain connectivity, one must include information on the chain correlations. For example, for Rouse chains, i j (r,r ) should be proportional to the pair correlators K i j defined in Eq. (3.34) [108, 241]. At first sight, using such complex Onsager coefficients in a simulation seems forbidding, but thanks to a clever trick due from Maurits and Fraaije [241], it becomes feasible [91, 242–245]. The DDFT method has been extended, e.g., to account for viscoelastic [246] or hydrodynamic effects [83, 247, 248]. A lattice version has also been proposed [249]. Time-Dependent Ginzburg-Landau (TDGL) Theories and Cell Dynamics Like dynamic density functional theories, time-dependent Ginzburg-Landau (TDGL) theories supplement a free energy functional F[ ] of an ‘order parameter’ field (r) by a model for the dynamical evolution of . The TDGL theories of interest in multiphase polymer systems mostly operate with locally conserved order parameters and have the same general structure as Eq. (3.63). For example, time-dependent Flory-Huggins de Gennes theories are used to study the demixing dynamics in polymer blends, and time-dependent Ohta-Kawasaki theories to study the ordering kinetics in copolymer systems. Discrete lattice versions of TDGL theories are often referred to as ‘cell dynamics’ models. We discuss specifically the time-dependent Ohta-Kawasaki theory. Starting from the free energy functional (3.42) for melts of diblock AB copolymers and choosing an Onsager coefficient that describes local diffusive dynamics, (r, r ) = M δ(r − r ), we obtain the dynamical equations [250] ∂ A δFFA + ∇(v A ) = M + η(r, t) ∂t δ A (r)  $ ρ0 #  ∂W ¯ A ) + η(r, t)  = − B A − A( A − N ∂ A

(3.65)

¯ A is the total volume fraction of monomers A in the system. Hence the some-what awkward long-range where ‘Coulomb’ term in Eq. (3.42) becomes short range, and the dynamical equations only depend on local terms. Because of the appealingly simple structure of the final theory, Eq. (3.65), it is widely used for simulations of copolymer systems at equilibrium and under shear (see below). Bahiana and Oono have formulated a discrete version on a lattice [251] which is equally popular in cell dynamics simulations. TDGL simulations are much faster than molecular field simulations, but of course, the underlying model is less accurate. Honda and Kawakatsu have recently proposed a multiscale hybrid method that combines the two approaches, using dynamic density functional input to improve on the accuracy of the TDGL model [167]. Such hybrid approaches will presumably gain importance in the future. 3.4.4

Applications

After this overview of the main models used for simulations of multiphase polymer systems, we will now illustrate them by reviewing simulation work that has been done on homopolymer blends and copolymer melts. We focus on simulations of generic models. Atomistic studies or studies of bottom-up models are scarce and have already been discussed earlier (Sections 3.4.1 and 3.4.2.1). 3.4.4.1

Homopolymer Blends

Bulk Properties We have already mentioned the pioneering simulations of binary blends by Sariban and Binder and by Cifra et al. [195–197] Following up on this work, a number of studies, mainly by Binder and coworkers, have considered the critical behavior of binary blends [44, 252–256]. As discussed in Section

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3.3.2.5, fluctuations shift the critical demixing point and change the critical exponents from mean-field-like to Ising-like. However, the region where the critical behavior deviates from the mean-field prediction shrinks with increasing chain length, and effective mean-field behavior could be observed already for relatively moderate chain lengths [252]. Coming from the other end, field-based Monte Carlo simulations of the FloryHuggins-de Gennes functional (3.13) have confirmed that noise shifts the coexistence curve and changes the critical behavior from mean-field to Ising [257]. Along with these studies of static blend properties, extensive work has also been dedicated to the dynamics of demixing. On the particle-based side, this was mainly investigated using kinetic Monte Carlo [242, 258, 259]. Reister et al. have compared results from kinetic Monte Carlo simulations and different versions of the dynamic density functional theory in a study of spinodal decomposition in symmetric blends [242]. Other field-based simulations have mainly relied on time-dependent Ginzburg Landau models, which have the advantage that one can reach much later stages of demixing. They were used to study demixing processes at equilibrium [126, 260–263] and under shear [264, 265]. Particularly interesting morphologies can be obtained if the dynamics in the two phases is distinctly different, i.e., one component becomes glassy during the demixing process [31] or crystallizes [29, 30, 266]. The particle-based studies mentioned so far have used coarse-grained models of blends that demix explicitly for energetic reasons. A number of authors have explored other factors that are believed to affect the chain miscibility with generic models [267], e.g., the effect of nonrandom mixing [268, 269], shape disparity [270], stiffness disparity [271], different architectures [272], and a different propensity towards crystallization [23, 27].

Internal Interfaces In the miscibility gap, polymer blends have highly inhomogeneous structures with droplets of one phase dispersed in the other phase. Their material properties are largely determined by the properties of the interfaces separating the two phases. While the distribution, size, and shape of the droplets depend on how the blend has been processed, the interfaces separating them often have time to reach local equilibrium and can be studied by means of equilibrium simulations. The first study of an interface in a binary blend was carried out by Reiter et al. [273] in 1990. Since then, several authors have investigated interfaces in symmetric [147, 148, 274–277] or asymmetric [278–280] binary blends (e.g., blends with stiffness and/or monomer size disparities) by means of generic particle-based simulations. The local interfacial structure is of interest because it determines the mechanical stability of an interface. For example, the local interfacial width gives the volume in which chains belonging to different phases can entangle. On the other hand, we have already discussed in Section 3.3.4 that interfaces exhibit capillary waves, which are significant on all length scales. This becomes apparent from the fact that the capillarywave contribution to the total width, as given by Eq. (3.53), diverges both if the system size L becomes very large and if the microscopic coarse-graining length a0 becomes very small. Therefore the length scales of the capillary waves and those of the local interfacial profiles cannot be separated clearly, and it is not clear, a priori, whether the concept of a ‘local interfacial structure’ is at all meaningful. This question has been investigated by Werner et al. [148] by simulations of the bond-fluctuation model (see Section 3.4.2.2). They demonstrated that it is indeed possible to describe homopolymer interfaces consistently in terms of a convolution of ‘intrinsic’ profiles with capillary wave undulations. They also studied the influence of confinement both on the capillary waves [281, 282] and on the intrinsic interfacial width [283]. Equilibrium simulations give information on the stability of interfaces under mechanical stress, but in a rather indirect way. A few authors have probed directly the rheological properties of interfaces with nonequilibrium particle- and field-based simulation methods, looking, e.g., at shear thinning and interfacial slip [284–287]. Detailed simulations of nonequilibrium interfaces are expensive, but with the development of new efficient simulation algorithms and modern fast computers, they become feasible.

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Surfaces and Films Another topic discussed intensely in the literature is the behavior of blends in the vicinity of surfaces, and that of blends confined to thin films. From a simulation point of view, these two situations are identical, because surfaces are typically studied in slab geometries (i.e., with periodic boundaries in two directions and free boundaries in the third). A large amount of work has been dedicated to the somewhat artificial situation of surfaces to which both types of monomers have exactly equal affinity. Even though these surfaces are perfectly neutral, one component will typically segregate to them: In incompatible blends, the minority component aggregates at the surface in order to reduce unfavorable contacts [288–290] (in contrast, the minority chains are removed from the surface in miscible blends [291]). In blends of chains with different stiffness, the stiffer chains are pushed towards the surface, because they loose less entropy there [292, 293]. For the same reason, linear chains aggregate at surfaces in blends of linear and star polymers [294]. Cavallo et al. have systematically investigated the phase behavior of films confined between neutral walls as a function of the film thickness. If the film is made thinner, one observes a crossover from three-dimensional to two-dimensional Ising behavior [295, 296]. Fluctuation effects in thin films are observed to be much stronger than in the bulk, consistent with our discussion in Section 3.3.2.5: In two dimensions, the Ginzburg parameter no longer scales with the chain length and stays finite for all chain lengths. A second transition occurs when the film becomes so thin that polymers are effectively two-dimensional, i.e., they can no longer pass each other. This reflects the theoretically expected fundamental difference between the demixing behavior of overlapping and non-overlapping two-dimensional polymers [15]. The more realistic situation of selective walls, to which one component adsorbs preferentially, has been addressed as well by different authors. In this case, the phase behavior is governed by wetting phenomena and capillary condensation [297–299]. The studies discussed so far were based on particle simulations. A few authors have used field-based simulations to explore dynamical aspects of phase separation in thin polymer blends. Morita et al. have studied the interplay of spinodal decomposition and interfacial roughening due to droplet formation with dynamic density functional simulations [300]. Shang et al. have used a time-dependent Ginzburg-Landau approach to study the spinodal phase separation of a thin film on a heterogeneous substrate [301].

3.4.4.2

Copolymer Systems

Copolymers as Compatibilizers Copolymers were originally designed as natural surfactant molecules that increase the miscibility of incompatible homopolymers and enhance their interfacial properties. They are usually much more expensive than their respective homopolymers, but adding a small amount of copolymer can already improve the properties of the homopolymer blend significantly. A number of researchers have considered the effect of copolymers on the demixing transition for different copolymer architectures [256, 302, 303]. Dadmun and Waldow [256] have pointed out that copolymers not only shift the transition point towards higher values of the Flory-Huggins parameter χ , but also change the critical exponents via a Fisher renormalization mechanism. In the phase-segregated regime, copolymers aggregate to the interface, reduce the interfacial tension, and enlarge the interfacial width. The interfacial structure of homopolymer interfaces with adsorbed copolymers has been explored in detail by several authors [277, 302, 304–307]. Milner and Xi noted in 1996 that the main compatibilizing effect of copolymers probably has a kinetic origin: [308] Copolymers reduce the rate of droplet coalescence during the processing of the blend via a Marangoni effect: If the copolymer concentration drops somewhere at the surface of a droplet, the surface tension increases locally. This induces flow in the direction of the weak point. Hence the copolymer film stabilizes itself kinetically, much like a soap film. In addition, the copolymer blocks stretching into the bulk

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form repulsive coronae. Experimental studies suggest that the resulting steric repulsion between droplets may be even more prohibitive for droplet coalescence than the Marangoni effect [309]. Kim and Jo have studied the influence of copolymers on the dynamics of demixing in a series of works [310–314], but they used kinetic Monte Carlo, hence they did not include hydrodynamic effects and could not study the effect of the Marangoni flow. If one increases the copolymer concentration beyond a certain threshold, the macrophase separated phase at high χ N eventually gives way to a microphase separated phase (see, e.g., the phase diagram shown in Figure 3.6). A few authors have explored the full phase behavior of ternary systems [110, 111, 315, 316]. In ternary systems containing A and C homopolymers and ABC triblock copolymers, a whole zoo of new tricontinuous gyroid phases can be observed [317]. Pure Bulk Copolymer Melts The propensity of copolymers to self-assemble into complex mesostructures makes them attractive for various micro- and nanotechnological applications, which is why pure copolymer melts have become interesting in their own right. Many simulation studies are now concerned with the properties of pure copolymer melts. Early studies of fluctuations and chain correlations near the order/disorder transition (ODT), coming from the disordered side [46, 47, 318, 319] have reproduced the ODT singularity in the structure factor and revealed the dumbbell structure of the chains mentioned earlier [47]. Later, intense work has been devoted to studying ordered lamellar structures below the ODT in melts of symmetrical diblock copolymers, and analyzing them with respect to their structure, their dynamics, and their fluctuations [320–331]. Simulation studies of asymmetric diblock copolymer melts have also reproduced most of the other mesophases – the cylindrical phase, the bcc sphere phase, and even the gyroid phase [204, 211, 235, 332–335]. Locating the actual position of the ODT accurately is difficult, especially in lattice models [336], since the natural periodicity of the structures is in general incompatible with the box size. This results in complex finite size artefacts. Nevertheless, phase diagrams have been obtained in recent years [337–339]. Particle-based and field-based simulations have provided evidence that for symmetrical diblock copolymers ( f = 1/2), the transition to the lamellar phase is shifted [224, 340], compared to the mean-field phase diagram, and becomes first order [340–342], in agreement with the theoretical expectation [116] (see Section 3.3.2.5). Two examples of phase diagrams obtained with different simulation methods are shown in Figure 3.8: One was calculated by Matsen et al. [338] using Monte Carlo simulations of copolymers with length N = 30 in a simple lattice model (left), and one by Lennon et al. [342] using field-theoretical calculations (right) at Ginzburg parameter C = 50. Both phase diagrams significantly improve on the SCF phase diagram of Figure 3.5 (right) in the weak segregation regime, and reproduce the main qualitative features of the experimental phase diagram (Figure 3.5, left): The transitions are first order everywhere. The ODT is shifted to higher χ N. Direct phase transition between the disordered phase and the complex mesophase (PL or G, respectively) or the lamellar phase are possible for a range of copolymer block fractions f . The phase diagrams even reproduce a small hump of the ODT transition line at the boundary to the complex mesophase (PL or G), which is also observed experimentally. The only ‘problem’ with the Monte Carlo phase diagram is that it features a perforated lamellar (PL) phase instead of a gyroid (G) phase. This may be a finite size artefact, as suggested by the authors, or a property of the lattice model under consideration. (Gyroid phases have been found in lattice models [204], but it should be noted that the free energies of the PL and the G phase are very close according to SCF calculations.) Nevertheless, the simulations demonstrate convincingly that the discrepancies between the experimental phase diagram and the SCF phase diagram shown in Figure 3.5 can to a large extent be attributed to the effect of fluctuations. In recent years, researchers have also begun to simulate melts of more complex copolymers, e.g., starblock copolymers [343–347], rod-coil copolymers [348, 349], diblock copolymers with one crystallizing component

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18

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C

C

L

50

16

PL

χeff N 14

χ eff N 25

12

disordered 0 0.0

0.2

0.4

0.6

f

0.8

1.0

10 0.3

0.4

0.5

f

Figure 3.8 Phase diagrams for diblock copolymer melts as obtained from Monte Carlo simulations of a lattice model (chain length N = 30) by Matsen et al. (left) [338] and from field-theoretical Complex-Langevin simulations (Ginzburg parameter C = 50) by Lennon et al. (right). [342] Symbols L,C,G correspond to diblock phases of Figure 3.3; in addition, the Monte Carlo phase diagram features a perforated lamellar (PL) phase. Left figure: Courtesy of reference [342]. Right figure: Reprinted from reference [342], Copyright (2008) with permission from American Physical Society.

[24–26, 350–352], triblock copolymers [353–356], or random copolymers [357, 358]. An interesting study on diblock copolymers with one amphiphilic block has recently been presented by Khokholov and Khalatur [359]. Since the amphiphilic block favors interfaces, the morphology of the mesophases changes completely and is characterized by thin channels and slits. Lay et al. and Palmer et al. have studied the computationally challenging problem of microphase separation in randomly crosslinked binary blends [360–362], and also the inverse problem, to which extent ordered copolymer structures can be stabilized by crosslinking [363]. A great deal of simulation work on copolymer melts has been done with timedependent Ginzburg-Landau approaches [364–378]. These studies have mostly addressed dynamical questions, i.e., the kinetics of ordering, disordering processes in pure melts [366–374] and in mixtures containing copolymers [375–378]. Already in pure diblock copolymer melts, ordering/disordering processes were found to proceed via intricate pathways that involve nontrivial intermediate states (e.g., the perforated lamellar state which is only metastable at equilibrium). In mixtures, the interplay of macrophase and microphase separation leads to a wealth of new transient morphologies [375–378] (see also Ref. 379). Confined Copolymer Melts In recent years, there has been growing interest in confined copolymer systems. The first studies have explored the ordering of copolymers melts in thin films between neutral walls [380–383] or general (selective) walls [384–394], both with particle-based models and dynamic density field theories. Triblock copolymers [395–400] and copolymer blends [382] have also been considered. The dynamics of copolymer ordering in confinement was studied with time-dependent Ginzburg-Landau methods [402]. A series of papers have dealt with the technologically relevant question, whether and how surface patterns can be transferred into copolymer films [389, 403–408]. The current focus shifts to nanocylindrically or spherically confined blends [205, 409–418]. A variety of new structures can already be observed in systems of diblock copolymers confined to nanocylinders, e.g., mesh structures, single and double helices [410]. For triblock copolymers, the spectrum is even more diverse [205]. The confined structures depend on the bulk structure and on the shape of the confining channels. Li and coworkers have shown that possible morphologies can be screened efficiently with the method of simulated annealing [419–424].

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Copolymers Under Shear Finally in this section, we briefly mention the large amount of work that has addressed the effect of shear on the microstructure of copolymer systems, with both field-based and particlebased methods [425–441]. Special attention has been given to the phenomenon that lamellae reorient under shear [431, 442–445]. Experimentally, it is observed that lamellae orient parallel to the shear flow at low shear rates, and perpendicular to the shear flow at high shear rates. Simulations give conflicting results. The parallel state consistently becomes unstable at high shear rates, but its relative stability at low shear rates (or results on the indicators of relative stability such as the entropy production) seems to depend on the type of model [443, 445].

3.5

Future Challenges

The equilibrium theory of fluid polymer mixtures is fairly advanced. Thanks to the universal properties of polymers, it requires relatively little input information (χ parameter etc.) to be predictive at a quantitative level. However, if one goes beyond equilibrium and beyond the fluid state, the situation is much less satisfying. For example, very little work has been done on crosslinked polymer blends [360–363], even though they are common in applications. Physical crosslinks can be established if domains crystallize or become glassy. As we have seen above, research on blends with one crystallizing or glassy component is also rather scarce. Another important issue is the influence of the blend processing on the properties of the resulting materials, i.e., the structure of phase separating blends under shear. The mechanical properties of immiscible blends depend crucially on their microscopic morphologies, i.e., the sizes and shapes of droplets, which in turn depend on the manufacturing process. Theoretical nonequilibrium state diagrams that relate the processing conditions (shear rates, geometry, copolymer concentration, etc.) with final morphologies are still missing. Since the relevant length scales are relatively large and hydrodynamics are important, the simulation method of choice should be a cell dynamics method that incorporates hydrodynamics, e.g., a Lattice-Boltzmann method [446]. Methods that combine Ginzburg-Landau functionals for immiscible fluids with Lattice-Boltzmann models for Newtonian fluids have been developed [447] and used to study demixing processes at rest [448] and under shear [449–451]. Giraud and coworkers have proposed a Lattice-Boltzmann method for viscoelastic fluids, which is more suitable to describe polymers [452], and carried out first simulations of viscoelastic liquid mixtures [453]. Nevertheless, the whole field is still in its infancy. As the field of polymer simulations reaches maturity, the bottom-up modeling approach (Section 3.4.1) will gain importance. So far, the vast majority of theoretical and simulation studies of (co)polymer blends was based on generic model systems. One or two decades from now, realistic simulations of specific polymer blends will probably be equally, if not more, common. One of the major challenges in this context is to develop hybrid multiscale methods that combine different levels of coarse-graining, i.e., use a relatively coarse basic model and fine-grain selectively in interesting regions of the materials (e.g., interfaces) [166].

Acknowledgements The author thanks Mark W. Matsen for introducing her to the self-consistent field theory a long time ago and for providing Figures 3.5 and 3.8 (left), and Glenn H. Fredrickson for discussions and for the permission to show Figure 3.8 (right). She has benefitted from collaborations and/or discussions with J¨org Baschnagel, Kurt Binder, Dominik D¨uchs, Burkhard D¨unweg, Avi Halperin, Venkat Ganesan, Kurt Kremer, Marcus M¨uller, Wolfgang Paul, Ulf Schiller, Jens Smiatek, Jens-Uwe Sommer, Andreas Werner, and many others. Financial support from the German Science Foundation (DFG) is gratefully acknowledged.

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References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44.

Flory P. J. Chem. Phys. 1941; 9: 660–661. Huggins M. J. Chem. Phys. 1941; 9: 440. Flory P. Statistical mechanics of chain molecules. John Wiley & Sons: New York, 1969. de Gennes P. Scaling concepts in polymer physics. Cornell University Press: Cornell, 1979. Doi M, Edwards S. The theory of polymer dynamics. Oxford University Press: New York, 1986. Fleer G, Cohen Stuart M, Scheutjens J, Cosgrove T, Vincent B. Polymers at interfaces. Kluwer Academic Publ: Reading, Massachusetts, 1993. Larson R. The structure and rheology of complex fluids. Oxford University Press: New York, 1999. Fredrickson G. The equilibrium theory of inhomogeneous polymers. Clarendon press: Oxford, 2006. Hamley I. The physics of block copolymers. Oxford University Press: Oxford, 1999. Guth E, Mark H. Monatshefte f¨ur Chemie 1934; 65: 93. Paul W, J B. Stochastic processes from physics to finance. Springer: Berlin, 1999. Kleinert H. Path integrals in quantum mechanics, polymer physics, and financial markets. 5th edn., World Scientific, 2009. Osa M, Yoshizaki T, Yamakawa H. Polymer JOURNAL 2004; 36: 634–641. Duplantier B, Saleur H. Nucl. Phys. B 1987; 290: 291–326. Semenov A, Johner A. Eur. Phys. J. E 2003; 12: 469–480. Zhulina EB, Borisov OV. Macromolecules 2008; 41: 5934–5944. Daoud M, Cotton J, Farnoux B, Jannink G, Sarma G, Benoit H, Duplessis R, Picot C, de Gennes P. Macromolecules 1975; 8: 804–818. Cacciuto A, Luijten E. Nano Lett. 2006; 6: 901–905. Wittmer J, Meyer H, Baschnagel J, Johner A, Obukhov S, Mattioni L, Muller M, Semenov A. Phys. Rev. Lett. 2004; 93: 147 801. Wittmer JP, Beckrich P, Johner A, Semenov AN, Obukhov SP, Meyer H, Baschnagel J. EPL 2007; 77: 56003. Ahlrichs P, Everaers R, D¨unweg B. Phys. Rev. E 2001; 64: 040 501. McLeish T. Adv. Phys. 2002; 51: 1379–1527. Hu W, Mathot V. J. Chem. Phys. 2003; 119: 10 953–10 957. Hu W. Macromolecules 2005; 38: 3977–3983. Hu W, Frenkel D. Faraday. Disc. 2005; 128: 253–260. Wang M, Hu W, Ma Y, Ma Y. J. Chem. Phys. 2006; 124: 244 901. Ma Y, Hu W, Wang H. Phys. Rev. E 2007; 76: 031 801. Ma Y, Zha L, Hu W, Reiter G, Han CC. Phys. Rev. E 2008; 77: 061 801. Mehta R, Kyu T. J. Pol. Sci. B 2004; 42: 2892–2899. Zhou D, Shi AC, Zhang P. J. Chem. Phys. 2008; 129: 154 901. Khalatur P, Khokholov A, Vasilevskaya V. Macrom. Theory and Sim. 1994; 3: 939–961. Genix A, Arbe A, Alvarez F, Colmenero J, Willner L, Richter D. Phys. Rev. E 2005; 72: 031 808. Han C, Bauer B, Clark J, Muroga Y, Matsushita Y, Okada M, Qui T, Chang T, Sanchez I. Polymer 1988; 29: 2002–2014. Hildebrand J, Scott R. The solubility of non-electrolytes. Dover: New York, 1964. Biros J, Zeman L, Patterson D. Macromolecules 1971; 4: 30–&. Freed K, Dudowicz J. Trends in Polymer Science 1995; 3: 248–255. Foreman K, Freed K. Adv. Chem. Phys. 1998; 103: 335–390. Dudowicz J, Freed K, Douglas J. Phys. Rev. Lett. 2002; 88: 095 503. Dudowicz J, Freed K, Douglas J. J. Chem. Phys. 2002; 116: 9983–9996. Freed K, Dudowicz J. Adv. Pol. Science 2005; 183: 63–126. Kumar S. Macromolecules 2000; 33: 5285–5291. Taylor J, Debenedetti P, Graessley W, Kumar S. Macromolecules 1996; 29: 764–773. Melenkevitz J, Crist B, Kumar S. Macromolecules 2000; 33: 6869–6877. M¨uller M, Binder K. Macromolecules 1995; 28: 1825–1834.

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Handbook of Multiphase Polymer Systems M¨uller M, Binder K. J. Physique II 1996; 6: 187–194. Fried H, Binder K. J. Chem. Phys. 1991; 94: 8349–8366. Fried H, Binder K. Europhys. Lett. 1991; 16: 237–242. Maurer W, Bates F, Lodge T, Almdal K, Mortensen K, Fredrickson G. J. Chem. Phys. 1998; 108: 2989–3000. Freed K, Dudowicz J. Macromolecules 1998; 31: 6681–6690. Dudowicz J, Freed K. Macromolecules 2000; 33: 3467–3477. Curro J, Schweizer K. J. Chem. Phys. 1988; 88: 7242–7243. Schweizer K. Macromolecules 1993; 26: 6033–6049. Schweizer K, Curro J. Adv. Chem. Phys. 1997; 98: 1–142. Wang Z. J. Chem. Phys. 2002; 117: 481–500. Akkermans RLC. J. Chem. Phys. 2008; 128: 244 904. Tambasco M, Lipson J, Higgins J. Macromolecules 2006; 39: 4860–4868. Lipson J. Macromolecules 1991; 24: 1334–1340. Lipson J, Andrews S. J. Chem. Phys. 1992; 96: 1426–1434. Tambasco M, Lipson J, Higgins J. Macromolecules 2004; 37: 9219–9230. Freed K, Pesci A. J. Chem. Phys. 1987; 87: 7342–7344. Shibayama M, Yang H, Stein R, Han C. Macromolecules 1985; 18: 2179–2187. Balsara N, Fetters L, Hadjichristidis N, Lohse D, Han C, Graessely W, Krishnamoorti R. Macromolecules 1992; 25: 6137–6147. Londono J, Wignall G. Macromolecules 1997; 30: 3821–3824. Fan C, Olafson B, Blanco M, Hsu S. Macromolecules 1992; 25: 3667–3676. Kipper M, Seifert S, Thiyagarajan P, Narasimhan B. Polymer 2004; 45: 3329–3340. Helfand E, Tagami Y. J. Pol. Sci. B Lett. 1971; 9: 741–&. Helfand E, Tagami Y. J. Chem. Phys. 1972; 57: 1812–&. Helfand E, Tagami Y. J. Chem. Phys. 1972; 56: 3592–&. Helfand E, Sapse A. J. Chem. Phys. 1975; 62: 1327–1331. Helfand E. J. Chem. Phys. 1975; 62: 999–1005. Schmid F. J. Phys.: Cond. Matt. 1998; 10: 8105–8138. Matsen M. J. Phys.: Cond. Matt. 2002; 14: R21–R47. FREED K. J. Chem. Phys. 1995; 103: 3230–3239. Hong K, Noolandi J. Macromolecules 1981; 14: 727–736. Schmid F. J. Chem. Phys. 1996; 104: 9191–9201. Matsen M, Schick M. Macromolecules 1994; 27: 6761–6767. Morse D, Fredrickson G. Phys. Rev. Lett. 1994; 73: 3235–3238. Duchs D, Sullivan D. J. Phys.: Cond. Matt. 2002; 14: 12 189–12 202. Hidalgo RC, Sullivan DE, Chen JZY. J. Phys.: Cond. Matt. 2007; 19: 376 107. Wang Q, Taniguchi T, Fredrickson G. J. Phys. Chem. B 2004; 108: 6733–6744. Cooke DM, Shi AC. Macromolecules 2006; 39: 6661–6671. Lynd NA, Hillmyer MA, Matsen MW. Macromolecules 2008; 41: 4531–4533. Hall DM, Lookman T, Fredrickson GH, Banerjee S. Phys. Rev. Lett. 2006; 97: 114 501. Maniadis P, Lookman T, Kober EM, Rasmussen KO. Phys. Rev. Lett. 2007; 99: 048 302. Feng EH, Lee WB, Fredrickson GH. Macromolecules 2007; 40: 693–702. Lee WB, Elliott R, Katsov K, Fredrickson GH. Macromolecules 2007; 40: 8445–8454. Thompson R, Ginzburg V, Matsen M, Balazs A. Science 2001; 292: 2469–2472. Reister E, Fredrickson G. J. Chem. Phys. 2005; 123: 214 903. Sides S, Kim B, Kramer E, Fredrickson G. Phys. Rev. Lett. 2006; 96: 250 601. Zinn-Justin J. Quantum field theory and critical phenomena. 2nd edn., Clarendon Press: Oxford, 1993. M¨uller M, Schmid F. Adv. Pol. Science 2005; 185: 1–58. Bates F, Schulz M, Khandpur A, Forster S, Rosedale J, Almdal K, Mortensen K. Faraday. Disc. 1994; 98: 7–18. F¨orster S, Khandpur A, Zhao J, Bates F, Hamley I, Ryan A, Bras W. Macromolecules 1994; 27: 6922–6935.

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94. Khandpur A, F¨orster S, Bates F, Hamley I, Ryan A, Bras W, Almdal K, Mortensen K. Macromolecules 1995; 28: 8796–8806. 95. Matsen M, Bates F. Macromolecules 1996; 29: 1091–1098. 96. Matsen M, Schick M. Phys. Rev. Lett. 1994; 72: 2660–2663. 97. Tyler C, Morse D. Phys. Rev. Lett. 2005; 94: 208 302. 98. Takenaka M, Wakada T, Akasaka S, Nishitsuji S, Saijo K, Shimizu H, Kim MI, Hasegawa H. Macromolecules 2007; 40: 4399–4402. 99. Scheutjens J, Fleer G. J. Phys. Chem. 1979; 83: 1619–1635. 100. Carignano M, Szleifer I. J. Chem. Phys. 1993; 98: 5006–5018. 101. Bryk P, MacDowell LG. J. Chem. Phys. 2008; 129: 104 901. 102. Tang H, Freed K. J. Chem. Phys. 1991; 94: 1572–1583. 103. Woodward C. J. Chem. Phys. 1991; 94: 3183–3191. 104. Woodward C, Yethiraj A. J. Chem. Phys. 1994; 100: 3181–3186. 105. Fraaije J. J. Chem. Phys. 1993; 99: 9202–9212. 106. Fraaije J, vanVlimmeren B, Maurits N, Postma M, Evers O, Hoffmann C, Altevogt P, GoldbeckWood G. J. Chem. Phys. 1997; 106: 4260–4269. 107. Hasegawa R, Doi M. Macromolecules 1997; 30: 3086–3089. 108. Kawakatsu T. Phys. Rev. E 1997; 56: 3240–3250. 109. Hillmyer M, Maurer W, Lodge T, Bates F, Almdal K. J. Phys. Chem. B 1999; 103: 4814–4824. 110. D¨uchs D, Ganesan V, Fredrickson G, Schmid F. Macromolecules 2003; 36: 9237–9248. 111. D¨uchs D, Schmid F. J. Chem. Phys. 2004; 121: 2798–2805. 112. Herktm¨atzki C, Schelten J. Phys. Rev. Lett. 1983; 51: 896–899. 113. Schwahn D, Mortensen K, Yeemadeira H. Phys. Rev. Lett. 1987; 58: 1544–1546. 114. Holyst R, Vilgis T. J. Chem. Phys. 1993; 99: 4835–4844. 115. Brazovskii S. Sov. Phys. JETP 1975; 41: 85. 116. Fredrickson G, Helfand E. J. Chem. Phys. 1987; 87: 697–705. 117. Shi A, Noolandi J, Desai R. Macromolecules 1996; 29: 6487–6504. 118. Leibler L. Macromolecules 1980; 13: 1602–1617. 119. Milner S, Olmsted P. J. Physique II 1997; 7: 249–255. 120. Erukhimovich I. Eur. Phys. J. E 2005; 18: 383–406. 121. Ohta T, Kawasaki K. Macromolecules 1986; 19: 2621–2632. 122. Ohta T, Kawasaki K. Macromolecules 1990; 23: 2413–2414. 123. Uneyama T, Doi M. Macromolecules 2005; 38: 196–205. 124. Yamada K, Nonomura M, Ohta T. J. Phys.: Cond. Matt. 2006; 18:L421–L427. 125. Kawakatsu T. Phys. Rev. E 1994; 50: 2856–2862. 126. Kielhorn L, Muthukumar M. J. Chem. Phys. 1999; 110: 4079–4089. 127. Sakurai J. Modern Quantum Mechanics. Addison Wesley, 1994. 128. Semenov A. Sov. Phys. JETP 1985; 61: 733–742. 129. Milner S, Witten T, Cates M. Europhys. Lett. 1988; 5: 413–418. 130. Zhulina E, Borisov O, Pryamitsyn V. J. Coll. Interf. Sci. 1990; 137: 495–511. 131. Zhulina E, Borisov O, Pryamitsyn V, Birshtein T. Macromolecules 1991; 24: 140–149. 132. Milner S, Witten T, Cates M. Macromolecules 1988; 21: 2610–2619. 133. Witten T, Leibler L, Pincus P. Macromolecules 1990; 23: 824–829. 134. Matsen M. J. Chem. Phys. 2001; 114: 10 528–10 530. 135. Netz R, Schick M. Europhys. Lett. 1997; 38: 37–42. 136. Netz R, Schick M. Macromolecules 1998; 31: 5105–5122. 137. Carelli C, Jones R, Young R, Cubitt R, Dalgliesh R, Schmid F, Sferrazza M. Phys. Rev. E 2005; 72: 031 807. 138. Fernandez M, Higgins J, Penfold J, Ward R, Shackleton C, Walsh D. Polymer 1988; 29: 1923–1928. 139. Stamm M, Schubert D. Ann. Rev. Mat. Sci. 1995; 25: 325–356. 140. Schubert D, Stamm M. Europhys. Lett. 1996; 35: 419–424. 141. Kerle T, Klein J, Binder K. Phys. Rev. Lett. 1996; 77: 1318–1321.

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142. Sferrazza M, Xiao C, Jones R, Bucknall D, Webster J, Penfold J. Phys. Rev. Lett. 1997; 78: 3693–3696. 143. Carelli C, Jones R, Young R, Cubitt R, Krastev R, Gutberlet T, Dalgliesh R, Schmid F, Sferrazza M. Europhys. Lett. 2005; 71: 763–769. 144. Buff F, Lovett R, Stillinger F. Phys. Rev. Lett. 1965; 15: 621–&. 145. Rowlinson J, Widom B. Molecular Theory of Capillarity. Clarendon: Oxford, 1982. 146. Semenov A. Macromolecules 1994; 27: 2732–2735. 147. Lacasse M, Grest G, Levine A. Phys. Rev. Lett. 1998; 80: 309–312. 148. Werner A, Schmid F, M¨uller M, Binder K. Phys. Rev. E 1999; 59: 728–738. 149. Schmid F, M¨uller M. Macromolecules 1995; 28: 8639–8645. 150. Forster D. Hydrodynamic fluctuations, broken symmetry, and correlation functions. Addison-Wesley: Reading, Massachusetts, 1995. 151. Baschnagel J, Binder K, Doruker P, Gusev A, Hahn O, Kremer K, Mattice W, M¨uller-Plathe F, Murat M, Paul W, et al.. Adv. Pol. Science 2000; 152: 41–156. 152. M¨uller-Plathe F. ChemPhysChem 2002; 3: 754–769. 153. M¨uller-Plathe F. Soft Materials 2003; 1: 1–31. 154. Binder K, Paul W, Sanots S, UW S. 2004; 152: 491–510. 155. Baschnagel J, Binder K, Paul W, Laso M, Suter U, Batoulis I, Jilge W, Burger T. J. Chem. Phys. 1991; 95: 6014–6025. 156. Paul W, Pistoor N. Macromolecules 1994; 27: 1249–1255. 157. Rapold R, Mattice W. J. Chem. Soc. – Far. Transactions 1995; 91: 2435–2441. 158. Carbone P, Varzaneh HAK, Chen X, Mueller-Plathe F. J. Chem. Phys. 2008; 128: 064 904. 159. Reith D, P¨utz M, M¨uller-Plathe F. J. Comp. Chem. 2003; 24: 1624–1636. 160. Faller R. Polymer 2004; 45: 3869–3876. 161. Milano G, M¨uller-Plathe F. J. Phys. Chem. B 2005; 109: 18 609–18 619. 162. Li X, Kou D, Rao S, Liang H. J. Chem. Phys. 2006; 124: 204 909. 163. Santangelo G, Di Matteo A, Muller-Plathe F, Milano G. J. Phys. Chem. B 2007; 111: 2765–2773. 164. Sewell TD, Rasmussen KO, Bedrov D, Smith GD, Thompson RB. J. Chem. Phys. 2007; 127: 144 901. 165. Guenza MG. J. Phys.: Cond. Matt. 2008; 20: 033 101. 166. Praprotnik M, Delle Site L, Kremer K. Ann. Rev. Phys. Chem. 2008; 59: 545–571. 167. Honda T, Kawakatsu T. Macromolecules 2007; 40: 1227–1237. 168. Akten E, Mattice W. Macromolecules 2001; 34: 3389–3395. 169. Rane S, Mattice W, Dhinojwala A. J. Phys. Chem. B 2004; 108: 14 830–14 839. 170. Choi P, Rane SS, Mattice WL. Macrom. Theory and Sim. 2006; 15: 563–572. 171. Sun Q, Faller R. J. Chem. Theory and Comp. 2006; 2: 607–615. 172. Sun Q, Faller R. J. Chem. Phys. 2007; 126: 144 908. 173. Sun Q, Pon FR, Faller R. Fluid Phase Equilibria 2007; 261: 35–40. 174. Choi K, Jo W. Macromolecules 1997; 30: 1509–1514. 175. Lee S, Lee J, Lee H, Mumby S. Polymer 1999; 40: 5137–5145. 176. Gestoso P, Brisson J. Polymer 2003; 44: 2321–2329. 177. Gestoso P, Brisson J. Polymer 2003; 44: 7765–7776. 178. Maranas J, Kumar S, Debenedetti P, Graessley W, Mondello M, Grest G. Macromolecules 1998; 31: 6998–7002. 179. Maranas J, Mondello M, Grest G, Kumar S, Debenedetti P, Graessley W. Macromolecules 1998; 31: 6991–6997. 180. Neelakantan A, Maranas J. Macromolecules 2004; 37: 8473–8480. 181. Faller R. Macromolecules 2004; 37: 1095–1101. 182. Rissanou AN, Peristeras LD, Economou IG. Polymer 2007; 48: 3883–3892. 183. Prathab B, Subramanian V, Aminabhavi TM. Polymer 2007; 48: 417–424. 184. Prathab B, Aminabhavi TM. Langmuir 2007; 23: 5439–5444. 185. Tao H, Fan C, Macknight W, Hsu S. Macromolecules 1994; 27: 1720–1728. 186. Spyriouni T, Vergelati C. Macromolecules 2001; 34: 5306–5316. 187. Yang H, Li Z, Lu Z, Sun C. Eur. Pol. J. 2005; 41: 2956–2962. 188. Jawalkar S, Adoor S, Sairam M, Nadagouda M, Aminabhavi T. J. Phys. Chem. B 2005; 109: 15 611–15 620.

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Jawalkar SS, Aminabhavi TM. Polymer 2006; 47: 8061–8071. Jawalkar SS, Raju KVSN, Halligudi SB, Sairam M, Aminabhavi TM. J. Phys. Chem. B 2007; 111: 2431–2439. Jawalkar SS, Nataraj SK, Raghu AV, Aminabhavi TM. J. Appl. Pol. Sci. 2008; 108: 3572–3576. Lee WJ, Ju SP, Wang YC, Chang JG. J. Chem. Phys. 2007; 127: 064 902. Lu Z, Li Z, Huang X, Jiang B, Sun J. Macrom. Theory and Sim. 1998; 7: 619–622. Pavel D, Shanks R. Polymer 2005; 46: 6135–6147. Sariban A, Binder K, Heermann D. Coll. Pol. Sci. 1987; 265: 424–431. Sariban A, Binder K. Macromolecules 1988; 21: 711–726. Cifra P, Karasz F, Macknight W. J. Pol. Sci. B 1988; 26: 2379–2383. Frenkel D, Smit B. Understanding molecular simulation. 2nd edn., Academic Press: San Diego, 2002. Ferrario M, Ciccotti G, Binder K ((eds.)). Computer simulations in condensed matter: From materials to chemical biology, vol. 2. Springer: Berlin, 2006. Larson R, Scriven L, Davis H. J. Chem. Phys. 1985; 83: 2411–2420. Dotera T, Hatano A. J. Chem. Phys. 1996; 105: 8413–8427. Pakula T. Macromolecules 1987; 20: 679–682. Carmesin I, Kremer K. Macromolecules 1988; 21: 2819–2823. Martinez-Veracoechea F, Escobedo F. Macromolecules 2005; 38: 8522–8531. Feng J, Ruckenstein E. J. Chem. Phys. 2007; 126: 124 902. Grest G, Lacasse M, Kremer K, Gupta A. J. Chem. Phys. 1996; 105: 10 583–10 594. Groot R, Warren P. J. Chem. Phys. 1997; 107: 4423–4435. Klapp S, Diestler D, Schoen M. J. Phys.: Cond. Matt. 2004; 16: 7331–7352. Laradji M, Guo H, Zuckermann M. Phys. Rev. E 1994; 49: 3199–3206. Soga K, Zuckermann M, Guo H. Macromolecules 1996; 29: 1998–2005. Besold G, Hassager O, Mouritsen O. Comp. Phys. Comm. 1999; 121: 542–544. Besold G, Guo H, Zuckermann M. J. Pol. Sci. B 2000; 38: 1053–1068. Detcheverry FA, Kang H, Daoulas KC, Mueller M, Nealey PF, de Pablo JJ. Macromolecules 2008; 41: 4989–5001. M¨uller M, Wittmer J, Cates M. Phys. Rev. E 1996; 53: 5063–5074. M¨uller M, Daoulas KC. J. Chem. Phys. 2008; 129: 164 906. Murat M, Kremer K. J. Chem. Phys. 1998; 108: 4340–4348. Eurich F, Maass P. J. Chem. Phys. 2001; 114: 7655–7668. Eurich F, Maass P, Baschnagel J. J. Chem. Phys. 2002; 117: 4564–4577. Eurich F, Karatchentsev A, Baschnagel J, Dieterich W, Maass P. J. Chem. Phys. 2007; 127: 134 905. Sambriski EJ, Guenza MG. Phys. Rev. E 2007; 76: 051 801. Fredrickson G, Ganesan V, Drolet F. Macromolecules 2002; 35: 16–39. Fredrickson GH. Soft Matter 2007; 3: 1329–1334. Lennon EM, Mohler GO, Ceniceros HD, Garcia-Cervera CJ, Fredrickson GH. Multiscale Modeling and Simulation 2007; 6: 1347–1370. Ganesan V, Fredrickson G. Europhys. Lett. 2001; 55: 814–820. Barrat J, Fredrickson G, Sides S. J. Phys. Chem. B 2005; 109: 6694–6700. Popov YO, Lee J, Fredrickson GH. J. Pol. Sci. B 2007; 45: 3223–3230. Lee J, Popov YO, Fredrickson GH. J. Chem. Phys. 2008; 128: 224 908. Moreira A, B¨aurle S, Fredrickson G. Phys. Rev. Lett. 2003; 91: 150 201. B¨aurle SA, Efimov GV, Nogovitsin EA. Europhys. Lett. 2006; 75: 378–384. B¨aurle S, Efimov G, Nogovitsin E. J. Chem. Phys. 2006; 124: 224 110. Risken H. The Fokker-Planck equation. 2nd edn., Springer: Berlin, 1989. Hoogerbrugge P, Koelman J. Europhys. Lett. 1992; 19: 155–160. Espanol P, Warren P. Europhys. Lett. 1995; 30: 191–196. Soddemann T, D¨unweg B, Kremer K. Phys. Rev. E 2003; 68: 046 702. Groot R, Madden T, Tildesley D. J. Chem. Phys. 1999; 110: 9739–9749. M¨uller M, Smith G. J. Pol. Sci. B 2005; 43: 934–958.

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Daoulas KC, M¨uller M. J. Chem. Phys. 2006; 125: 184 904. Ganesan V, Pryamitsyn V. J. Chem. Phys. 2003; 118: 4345–4348. Baeurle SA, Usami T, Gusev AA. Polymer 2006; 47: 8604–8617. Maurits N, vanVlimmeren B, Fraaije J. Phys. Rev. E 1997; 56: 816–825. Maurits N, Fraaije J. J. Chem. Phys. 1997; 107: 5879–5889. Reister E, M¨uller M, Binder K. Phys. Rev. E 2001; 64: 041 804. He X, Schmid F. Macromolecules 2006; 39: 2654–2662. He X, Schmid F. Macromolecules 2006; 39: 8908–8910. He X, Schmid F. Phys. Rev. Lett. 2008; 100: 137 802. Maurits N, Zvelindovsky A, Fraaije J. J. Chem. Phys. 1998; 109: 11 032–11 042. Maurits N, Zvelindovsky A, Sevink G, van Vlimmeren B, Fraaije J. J. Chem. Phys. 1998; 108: 9150–9154. Honda T, Kawakatsu T. J. Chem. Phys. 2008; 129: 114904. Mihajlovic M, Lo T, Shnidman Y. Phys. Rev. E 2005; 72: 041 801. Oono Y, Shiwa Y. Mod. Phys. Lett. B 1987; 1: 49–55. Bahiana M, Oono Y. Phys. Rev. A 1990; 41: 6763–6771. Deutsch H, Binder K. Macromolecules 1992; 25: 6214–6230. Deutsch H, Binder K. J. Physique II 1993; 3: 1049–1073. Binder K. Adv. Pol. Science 1994; 112: 181–299. M¨uller M, Wilding N. Phys. Rev. E 1995; 51: 2079–2089. Dadmun M, Waldow D. Phys. Rev. E 1999; 60: 4545–4550. Forrest B, Toral R. J. Stat. Phys. 1994; 77: 473–489. Sariban A, Binder K. Macromolecules 1991; 24: 578–592. Brown G, Chakrabarti A. Phys. Rev. E 1993; 48: 3705–3711. Takenaka M, Hashimoto T, Dobashi T. Phys. Rev. E 1995; 52: 5142–5145. Cao Y, Zhang H, Xiong Z, Yang Y. Macrom. Theory and Sim. 2001; 10: 314–324. Cheng M, Nauman E. J. Pol. Sci. B 2004; 42: 603–612. Prusty M, Keestra BJ, Goossens JGP, Anderson PD. Chem. Eng. Science 2007; 62: 1825–1837. He D, Nauman E. Chem. Eng. Science 1997; 52: 481–496. Zhang Z, Zhang H, Yang Y, Vinckier I, Laun H. Macromolecules 2001; 34: 1416–1429. Xu H, Chiu HW, Okabe Y, Kyu T. Phys. Rev. E 2006; 74: 011801. M¨uller M. Macrom. Theory and Sim. 1999; 8: 343–374. Kumar S. Macromolecules 1994; 27: 260–271. Yang Y, Lu J, Yan D, Ding J. Macrom. Theory and Sim. 1994; 3: 731–741. M¨uller M. Macromolecules 1995; 28: 6556–6564. Weinhold J, Kumar S, Singh C, Schweizer K. J. Chem. Phys. 1995; 103: 9460–9474. Theodorakis PE, Avgeropoulos A, Freire JJ, Kosmas M, Vlahos C. J. Phys.: Cond. Matt. 2007; 19: 466 111. Reiter J, Zifferer G, Olaj O. Macromolecules 1990; 23: 224–228. M¨uller M, Binder K, Oed W. J. Chem. Soc. - Far. Transactions 1995; 91: 2369–2379. Ypma G, Cifra P, Nies E, vanBergen A. Macromolecules 1996; 29: 1252–1259. Olaj O, Petrik T, Zifferer G. J. Chem. Phys. 1998; 109: 4646–4650. Qian H, Lu Z, Chen L, Li Z, Sun C. J. Chem. Phys. 2005; 122: 184 907. M¨uller M, Werner A. J. Chem. Phys. 1997; 107: 10 764–10 776. Adhikari N, Auhl R, Straube E. Macrom. Theory and Sim. 2002; 11: 315–325. Adhikari N, Straube E. Macrom. Theory and Sim. 2003; 12: 499–507. Werner A, Schmid F, M¨uller M, Binder K. J. Chem. Phys. 1997; 107: 8175–8188. Werner A, M¨uller M, Schmid F, Binder K. J. Chem. Phys. 1999; 110: 1221–1229. Binder K, M¨uller M, Schmid F, Werner A. J. Stat. Phys. 1999; 95: 1045–1068. Barsky S, Robbins M. Phys. Rev. E 2001; 63: 021 801. Barsky S, Robbins M. Phys. Rev. E 2002; 65: 021 808. Narayanan B, Pryamitsyn V, Ganesan V. Macromolecules 2004; 37: 10 180–10 194. Lo T, Mihajlovic M, Shnidman Y, Li W, Gersappe D. Phys. Rev. E 2005; 72: 040 801.

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Kumar S, Tang H, Szleifer I. Mol. Phys. 1994; 81: 867–872. Rouault Y, Baschnagel J, Binder K. J. Stat. Phys. 1995; 80: 1009–1031. Rouault Y, Dunweg B, Baschnagel J, Binder K. Polymer 1996; 37: 297–304. Cifra P, Karasz F, Macknight W. Macromolecules 1992; 25: 4895–4901. Yethiraj A, Kumar S, Hariharan A, Schweizer K. J. Chem. Phys. 1994; 100: 4691–4694. Kumar S, Yethiraj A, Schweizer K, Leermakers F. J. Chem. Phys. 1995; 103: 10 332–10 346. Batman R, Gujrati PD. J. Chem. Phys. 2008; 128: 124 903. Cavallo A, M¨uller M, Binder K. Europhys. Lett. 2003; 61: 214–220. Cavallo A, M¨uller M, Binder K. J. Phys. Chem. B 2005; 109: 6544–6552. Pereira G, Wang J. J. Chem. Phys. 1996; 104: 5294–5303. M¨uller M, Binder K. Macromolecules 1998; 31: 8323–8346. M¨uller M, Binder K. Phys. Rev. E 2001; 63: 021 602. Morita H, Kawakatsu T, Doi M. Macromolecules 2001; 34: 8777–8783. Shang Y, Kazmer D, Wei M, Mead J, Barry C. J. Chem. Phys. 2008; 128: 224 909. Dadmun M. Macromolecules 1996; 29: 3868–3874. Dai B, Song M, Hourston D, He X, Liang H, Pan C. Polymer 2004; 45: 1019–1026. Werner A, Schmid F, Binder K, M¨uller M. Macromolecules 1996; 29: 8241–8248. Werner A, Schmid F, M¨uller M. J. Chem. Phys. 1999; 110: 5370–5379. Liang H. Macromolecules 1999; 32: 8204–8209. Kim S, Jo W. J. Chem. Phys. 1999; 110: 12 193–12 201. Milner S, Xi H. J. Rheol. 1996; 40: 663–687. Lyu S, Jones T, Bates F, Macosko C. Macromolecules 2002; 35: 7845–7855. Jo W, Kim S. Macromolecules 1996; 29: 7204–7211. Kim S, Jo W, Kim J. Macromolecules 1996; 29: 6933–6940. Kim S, Jo W, Kim J. Macromolecules 1997; 30: 3910–3915. Kim S, Jo W. J. Chem. Phys. 1998; 108: 4267–4281. Ko M, Kim S, Jo W. Polymer 2000; 41: 6387–6394. M¨uller M, Schick M. J. Chem. Phys. 1996; 105: 8885–8901. Poncela A, Rubio A, Freire J. J. Chem. Phys. 2003; 118: 425–433. Dotera T. Phys. Rev. Lett. 2002; 89: 205 502. Binder K, Fried H. Macromolecules 1993; 26: 6878–6883. Pakula T, Karatasos K, Anastasiadis S, Fytas G. Macromolecules 1997; 30: 8463–8472. Weyersberg A, Vilgis T. Phys. Rev. E 1993; 48: 377–390. Larson R. Macromolecules 1994; 27: 4198–4203. Larson R. Macromolecules 1995; 28: 4784. Pan X, Shaffer J. Macromolecules 1996; 29: 4453–4455. Hoffmann A, Sommer J, Blumen A. J. Chem. Phys. 1997; 106: 6709–6721. Banaszak M, Clarke J. Phys. Rev. E 1999; 60: 5753–5756. Jo W, Jang S. J. Chem. Phys. 1999; 111: 1712–1720. Murat M, Grest G, Kremer K. Europhys. Lett. 1998; 42: 401–406. Murat M, Grest G, Kremer K. Macromolecules 1999; 32: 595–609. Wang Q, Nealey P, de Pablo J. Macromolecules 2002; 35: 9563–9573. Yang J, Winnik M, Pakula T. Macrom. Theory and Sim. 2005; 14: 9–20. Srinivas G, Swope WC, Pitera JW. J. Phys. Chem. B 2007; 111: 13 734–13 742. Hoffmann A, Sommer J, Blumen A. J. Chem. Phys. 1997; 107: 7559–7570. Groot R, Madden T. J. Chem. Phys. 1998; 108: 8713–8724. Horsch M, Zhang Z, Iacovella C, Glotzer S. J. Chem. Phys. 2004; 121: 11 455–11 462. Martinez-Veracoechea FJ, Escobedo FA. J. Chem. Phys. 2006; 125: 104 907. Micka U, Binder K. Macrom. Theory and Sim. 1995; 4: 419–447. Schultz A, Hall C, Genzer J. J. Chem. Phys. 2002; 117: 10 329–10 338. Matsen M, Griffiths G, Wickham R, Vassiliev O. J. Chem. Phys. 2006; 124: 024 904.

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Beardsley TM, Matsen MW. Eur. Phys. J. E 2008; 27: 323–333. Vassiliev O, Matsen M. J. Chem. Phys. 2003; 118: 7700–7713. M¨uller M, Daoulas KC. J. Chem. Phys. 2008; 128: 024 903. Lennon EM, Katsov K, Fredrickson GH. Phys. Rev. Lett. 2008; 101: 138 302. Gemma T, Hatano A, Dotera T. Macromolecules 2002; 35: 3225–3237. Xu Y, Feng J, Liu H, Hu Y. 2006; 32: 375–383. Xu Y, Feng J, Liu H, Hu Y. 2008; 34: 559–565. He X, Huang L, Liang H, Pan C. J. Chem. Phys. 2002; 116: 10 508–10 513. He X, Huang L, Liang H, Pan C. J. Chem. Phys. 2003; 118: 9861–9863. AlSunaidi A, Den Otter W, Clarke J. Phil. Trans. A 2004; 362: 1773–1781. Pryamitsyn V, Ganesan V. J. Chem. Phys. 2004; 120: 5824–5838. Hu W, Karssenberg FG, Mathot VBF. Polymer 2006; 47: 5582–5587. Wang M, Hu W, Ma Y, Ma Y. J. Chem. Phys. 2006; 124: 244 901. Qian Y, Cai T, Hu W. Macromolecules 2008; 41: 7625–7629. Ko M, Kim S, Jo W. Macrom. Theory and Sim. 2001; 10: 381–388. Banaszak M, Woloszczuk S, Pakula T, Jurga S. Phys. Rev. E 2002; 66: 031 804. Huh J, Jo W, ten Brinke G. Macromolecules 2002; 35: 2413–2416. Song J, Shi T, Li Y, Chen J, An L. J. Chem. Phys. 2008; 129: 054 906. Houdayer J, M¨uller M. Europhys. Lett. 2002; 58: 660–665. Houdayer J, M¨uller M. Macromolecules 2004; 37: 4283–4295. Khokhlov AR, Khalatur PG. Chem. Phys. Lett. 2008; 461: 58–63. Lay S, Sommer J, Blumen A. J. Chem. Phys. 2000; 113: 11 355–11 363. Lay S, Sommer J, Blumen A. J. Chem. Phys. 2000; 113: 11 355–11 363. Palmer K, Lastoskie C. J. Chem. Phys. 2004; 120: 486–495. Lay S, Sommer J, Blumen A. J. Chem. Phys. 1999; 110: 12173–12182. Oono Y, Bahiana M. Phys. Rev. Lett. 1988; 61: 1109–1111. Chakrabarti A, Toral R, Gunton J. Phys. Rev. A 1991; 44: 6503–6507. Shiwa Y, Taneike T, Yokojima Y. Phys. Rev. Lett. 1996; 77: 4378–4381. Komura S, Fukuda J, Paquette G. Phys. Rev. E 1996; 53: R5588–R5591. Qi S, Wang Z. Phys. Rev. Lett. 1996; 76: 1679–1682. Qi S, Wang Z. Phys. Rev. E 1997; 55: 1682–1697. Zhang H, Zhang J, Yang Y, Zhou X. J. Chem. Phys. 1997; 106: 784–792. Hamley I. Macrom. Theory and Sim. 2000; 9: 363–380. Hamley I, Stanford J, Wilkinson A, Elwell M, Ryan A. Polymer 2000; 41: 2569–2576. Ren S, Hamley I. Macromolecules 2001; 34: 116–126. Vega D, Harrison C, Angelescu D, Trawick M, Huse D, Chaikin P, Register R. Phys. Rev. E 2005; 71: 061 803. Ohta T, Ito A. Phys. Rev. E 1995; 52: 5250–5260. Ohta T, Motoyama M, Ito A. J. Phys.: Cond. Matt. 1996; 8: A65–A80. Ito A. Phys. Rev. E 1998; 58: 6158–6165. Zhang J, Jin G, Ma Y. J. Phys.: Cond. Matt. 2006; 18: 837–845. Morita H, Kawakatsu T, Doi M, Yamaguchi D, Takenaka M, Hashimoto T. Macromolecules 2002; 35: 7473–7480. Sommer J, Hoffmann A, Blumen A. J. Chem. Phys. 1999; 111: 3728–3732. Sommer J, Hoffmann A, Blumen A. J. Chem. Phys. 1999; 111: 3728–3732. Song J, Li Y, Huang Q, Shi T, An L. J. Chem. Phys. 2007; 127: 094 903. Alexander-Katz A, Fredrickson GH. Macromolecules 2007; 40: 4075–4087. Kikuchi M, Binder K. Europhys. Lett. 1993; 21: 427–432. Kikuchi M, Binder K. J. Chem. Phys. 1994; 101: 3367–3377. Geisinger T, M¨uller M, Binder K. J. Chem. Phys. 1999; 111: 5241–5250. Geisinger T, M¨uller M, Binder K. J. Chem. Phys. 1999; 111: 5251–5258. Huinink H, Brokken-Zijp J, van Dijk M, Sevink G. J. Chem. Phys. 2000; 112: 2452–2462.

380. 381. 382. 383. 384. 385. 386. 387. 388.

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Wang Q, Yan Q, Nealey P, de Pablo J. J. Chem. Phys. 2000; 112: 450–464. Feng J, Ruckenstein E. Macrom. Theory and Sim. 2002; 11: 630–639. Pereira G. Eur. Phys. J. E 2002; 7: 273–289. Feng J, Liu H, Hu Y, Prausnitz J. Chem. Eng. Science 2004; 59: 1701–1710. Knoll A, Lyakhova K, Horvat A, Krausch G, Sevink G, Zvelindovsky A, Magerle R. Nature Materials 2004; 3: 886–890. Lyakhova K, Sevink G, Zvelindovsky A, Horvat A, Magerle R. J. Chem. Phys. 2004; 120: 1127–1137. Szamel G, M¨uller M. J. Chem. Phys. 2003; 118: 905–913. Ludwigs S, B¨oker A, Voronov A, Rehse N, Magerle R, Krausch G. Nature Materials 2003; 2: 744–747. Horvat A, Lyakhova K, Sevink G, Zvelindovsky A, Magerle R. J. Chem. Phys. 2004; 120: 1117–1126. Lyakhova K, Horvat A, Zvelindovsky A, Sevink G. Langmuir 2006; 22: 5848–5855. Huang Y, Liu H, Hu Y. Macrom. Theory and Sim. 2006; 15: 117–127. Xiao X, Huang Y, Liu H, Hu Y. 2007; 33: 1083–1091. Brown G, Chakrabarti A. J. Chem. Phys. 1994; 101: 3310–3317. Brown G, Chakrabarti A. J. Chem. Phys. 1995; 102: 1440–1448. Pereira G, Williams D, Chakrabarti A. J. Chem. Phys. 2000; 112: 10 011–10 016. Wang Q, Nath S, Graham M, Nealey P, de Pablo J. J. Chem. Phys. 2000; 112: 9996–10 010. Jayaraman A, Hall C, Genzer J. J. Chem. Phys. 2005; 123: 124 702. Daoulas K, M¨uller M, Stoykovich M, Park S, Papakonstantopoulos Y, de Pablo J, Nealey P, Solak H. Phys. Rev. Lett. 2006; 96: 036 104. Daoulas KC, M¨uller M, Stoykovich MP, Kang H, de Pablo JJ, Nealey PF. Langmuir 2008; 24: 1284–1295. Chen H, Peng C, Sun L, Liu H, Hu Y, Jiang J. Langmuir 2007; 23: 11 112–11 119. Xu JB, Wu H, Lu DY, He XF, Zhao YH, Wen H. 2006; 32: 357–362. Feng J, Ruckenstein E. Macromolecules 2006; 39: 4899–4906. Feng J, Ruckenstein E. J. Chem. Phys. 2006; 125: 164 911. Feng J, Liu H, Hu Y. Macrom. Theory and Sim. 2006; 15: 674–685. Feng J, Liu H, Hu Y. Fluid Phase Equilibria 2007; 261: 50–57. Feng J, Liu H, Hu Y, Jiang J. Macrom. Theory and Sim. 2008; 17: 163–170. Feng J, Ruckenstein E. J. Chem. Phys. 2008; 128: 074 903. Wang Q. J. Chem. Phys. 2007; 126: 024 903. Xiao X, Huang Y, Liu H, Hu Y. Macrom. Theory and Sim. 2007; 16: 732–741. Han Y, Cui J, Jiang W. Macromolecules 2008; 41: 6239–6245. Yin Y, Sun P, Chen T, Li B, Jin Q, Ding D, Shi A. ChemPhysChem 2004; 5: 540–548. Yin Y, Sun P, Jiang R, Li B, Chen T, Jin Q, Ding D, Shi C. J. Chem. Phys. 2006; 124: 184 708. Yu B, Sun P, Chen T, Jin Q, Ding D, Li B, Shi AC. J. Chem. Phys. 2007; 126: 204 903. Yu B, Sun P, Chen T, Jin Q, Ding D, Li B, Shi AC. J. Chem. Phys. 2007; 127: 114 906. Yu B, Li B, Jin Q, Ding D, Shi AC. Macromolecules 2007; 40: 9133–9142. Chen P, Liang H. J. Phys. Chem. B 2008; 112: 1918–1925. Ohta T, Enomoto Y, Harden J, Doi M. Macromolecules 1993; 26: 4928–4934. Kodama H, Doi M. Macromolecules 1996; 29: 2652–2658. Zvelindovsky A, van Vlimmeren B, Sevink G, Maurits N, Fraaije J. J. Chem. Phys. 1998; 109: 8751–8754. Zvelindovsky A, Sevink G, van Vlimmeren B, Maurits N, Fraaije J. Phys. Rev. E 1998; 57: R4879–R4882. Zhang K, Manke C. Comp. Phys. Comm. 2000; 129: 275–281. Hamley I. J. Phys.: Cond. Matt. 2001; 13: R643–R671. Ren S, Hamley I, Sevink G, Zvelindovsky A, Fraaije J. Macrom. Theory and Sim. 2002; 11: 123–127. Sevink G, Zvelindovsky A. J. Chem. Phys. 2004; 121: 3864–3873. Feng J, Ruckenstein E. J. Chem. Phys. 2004; 121: 1609–1625. Honda T, Kawakatsu T. Macromolecules 2006; 39: 2340–2349. Rychkov I. Macrom. Theory and Sim. 2005; 14: 207–242. Liu D, Zhong C. Macrom. Rap. Comm. 2005; 26: 1960–1964. Narayanan B, Pryamitsyn V, Ganesan V. Phys. Rev. Lett. 2006; 96: 028 302.

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Li X, Tang P, Zhang H, Qiu F, Yang Y. J. Chem. Phys. 2008; 128: 114 901. You LY, Chen LJ, Qian HJ, Lu ZY. Macromolecules 2007; 40: 5222–5227. Pinna M, Zvelindovsky AV, Todd S, Goldbeck-Wood G. J. Chem. Phys. 2006; 125: 154 905. Pinna M, Zvelindovsky AV. Soft Matter 2008; 4: 316–327. Fraser B, Denniston C, Muser M. J. Pol. Sci. B 2005; 43: 970–982. Fraser B, Denniston C, Muser M. J. Chem. Phys. 2006; 124: 104 902. Liu W, Qian HJ, Lu ZY, Li ZS, Sun CC. Phys. Rev. E 2006; 74: 021 802. Lisal M, Brennan JK. Langmuir 2007; 23: 4809–4818. Succi S. The Lattice Boltzmann equation for fluid dynamics and beyond. Oxford University Press, 2001. Orlandini E, Swift M, Yeomans J. Europhys. Lett. 1995; 32: 463–468. Wagner A, Yeomans J. Phys. Rev. Lett. 1998; 80: 1429–1432. Wagner A, Yeomans J. Int. J. Mod. Phys. C 1997; 8: 773–782. Wagner A, Yeomans J. Phys. Rev. E 1999; 59: 4366–4373. Martys N, Douglas J. Phys. Rev. E 2001; 63: 031 205. Giraud L, D’Humieres D, Lallemand P. Europhys. Lett. 1998; 42: 625–630. Wagner A, Giraud L, Scott C. Comp. Phys. Comm. 2000; 129: 227–232. Morkved TL, et al. Dynamics of ternary polymer blends: Disordered, ordered and bicontinuous microemulsion phases. Faraday Discussions/112/0301-7239, pp 339, 1999.

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4 Interfaces in Multiphase Polymer Systems Gy¨orgy J. Marosi Budapest University of Technology and Economics, Budapest, Hungary

4.1 Introduction Interfaces, in their widest meaning, divide the space into distinct compartments, thus the interface is one of the basic phenomena of the organization of the world. The segregation of certain organic molecules into multiphase systems, assisted by phase borders, was an essential condition for the origin of life. Concerning multiphase polymeric materials, their behavior is determined, beyond the bulk characteristics of the phases and their ratio, by the intermediate phase (interphase) of altered structure and molecular dynamics. The role of material located in the interfacial region is far larger in the determination of the macroscopic characteristics of multiphase systems than its share in volume or mass. Refining the phase structure increases the share and importance of the interfacial region and this growth accelerates significantly in the submicron range. In nanostructured multiphase polymer systems, in which the specific surface area of phase borders can be several orders of magnitude larger than between microphase structures, almost all the polymer chains belong to the interphase (the bulk disappears). The extent of the influence of the phase contact on their structure and mobility depends on the distance and type of interaction. The interfacial interactions determine the formation of the structure at the phase borders and consequently its physical characteristics, chemical stability and biological effects. The interface-related physical effects include alteration of strength, transport and electrical characteristics. The stability of (nano)composites is influenced not only by physical effects (such as interfacial heat conductivity) but also by chemical effects of catalytic ions and molecular mobility at the interfaces. The structure and mobility of biomolecules at phase borders influence the bioavailability of drug delivery systems, the immune response around implants, the strength of biomaterials and the controllability of all biological effects. As the interface-related fields of application include the use of (nano)composites, electrical units, packaging materials, stabilizing agents, catalysts, tribological agents, and biomaterials the relevant market shear is huge. The importance of nanocomposites and nanomedicine is based basically on their enhanced interfacial effects. Interface design is relevant even in the waste management, therefore a wide scope has to be considered when the interfacial phenomena of the multiphase systems are discussed.

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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In the past, the separate discussions about the different interface-related fields, including the artificial distinction between flat and particulate surfaces/interfaces, limited the synergistic common initiatives between the representatives of the relevant disciplines. The aim of this chapter is to show the common fundamentals, similar approaches and convergent activities in the different fields of multiphase polymer and biopolymer systems and to try to promote the common thinking of these research communities. The first step for achieving this aim was to establish common classifications concerning characteristics, modification and analyses of the interfaces (including both flat and particulate ones). The evaluation of thermodynamic and kinetic factors establishes a common basis for all types of modifications. The relationship between the macroscopic properties and the interfacial structure has been most widely studied in the case of mechanical response, which is also reviewed briefly. The most challenging task is to categorize the different types of interfacial modifications performed for various purposes and industrial segments. Modifications of morphology, segmental mobility, compatibility, reactivity and responsiveness has been considered, but still some important areas such as catalysis and tribology are not involved in the scope of this study.

4.2

Basic Considerations

Some relevant terms have to be summarized before discussing the main interfacial characteristics. The term surface means a border between phases of different state matter, while the term interface is a sharp boundary between phases of the same (e.g. solid–solid) state (the difference is artificial as both of them can be considered as interface). If a layer exists between the phases, differing from the bulk materials, it is called the interphase (IP), interlayer, or interfacial zone. The interphases form boundary layers of different width usually in the micro- or nanometer range. The use of the terms ‘interphase’ or ‘interface’ depends in most cases on the point of view (micro- or nanometer scale). The extent and nature of the IP hinge upon how the interphase formation, governed by both thermodynamic and kinetic effects, takes place. On the micro-scale, the IP can be considered a distinct 3D continuum of average mechanical and other properties. On the nano-scale the relaxation processes of polymer segments in the region of interfacial influence has to be considered. Increasing deviation from Gaussian chain statistics occurs as the effect of interface becomes dominant. A ‘train-loop-tail’ structure of restricted mobility is characteristic of the chains near to a solid surface (excepting the particular case of polymer brushes grown perpendicularly to the surface) in contrast to the random distribution of the coil of mobile chains in a melt (see in Figure 4.1).

Figure 4.1 Interfacial structures.

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INTERFACE characteristics

structural charge

mobility

morphological geometrical

conducting amorphous semicond. crystalline insulating oriented

volume thickness hierarchy

interaction

migration segmental sorption interfacial bonds crosslinking degradation barrier gateway

restricted enhanced varied

ad~ de~

secondary acid-base covalent

activation inhibition

activation inhibition

Figure 4.2 Classification of main characteristics related to interfaces.

As a good interfacial contact restricts the number of possible conformations available for the chains, their entropy is reduced. The internal energy of chains can be altered too (depending on the energy of the surface–polymer interaction). In nanocomposites the average interparticle distance can be reduced to the magnitude of the radii of gyration (Rg ) of the chains (at a few % nano-particle content) and all the chains of the system become more or less immobilized. In these systems the IP means the entropic difference among the entangled segments at different level of immobilization depending on their distance from the phase border. The thickness of IP cannot be defined exactly as the level of immobilization changes continuously. The IP zone is formed in spontaneous process or engineered consciously when phases of different character (such as hard–soft, charged–neutral, polar–apolar, living–lifeless) are attached to each other. The rules governing the phenomena (and determining the design of IPs) are basically the same in all kinds of materials. Apart from certain differences, various similarities can be found in the surface-related aspects of biocompatibility, tribology, packaging, blending, fire protection, catalysis, formation of (nano-)composites, optoelectronical devices, and other multiphase systems.

4.3 Characteristics of Interfacial Layers The characteristics of interfaces and interfacial interactions are summarized in Figure 4.2. Structural characteristics include the molecular/segmental and charge arrangement in material units taking part in the formation of micro- or nanoscale interlayers. The morphological features of the interfacial region are influenced by the interaction between the associated groups of the phases, mobility of adjacent segments and local shear forces. Strong interaction, that restricts the molecular or segmental mobility, leads to stabile amorphous structure in the interfacial region (requested in many pharmaceuticals). The interactions, however, can also transmit a template for certain crystalline organization, while local shear forces initiate orientation (favored in many polymer composites). The volume of the interlayers within the multiphase system increases by several orders of magnitude when the size of inclusions is reduced from micro- to nanoscale. Good wetting of a rough surface is also accompanied with the increase of the internal contact area between the chains and the inclusions. The thickness of the IP, falling in the range of thin Langmuir−Blodgett (LB) monomolecular films and macroscopic layer of several hundred μm, is determined by the balance between entropy (segmental arrangement) and free energy (adhesion forces) effects. Hierarchic structure of layers of different thickness is characteristic to natural multiphase structures and to some advanced composites [1]. An interface separates and/or integrates the phases and a separating interlayer (containing for example inorganic nanoparticles) can act also as a joining site between incompatible phases.

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Figure 4.3 Fibers embedded in an interpenetrating network of A () polymer (adhering to the fiber) and B (≡) polymer (non-adhering to the fiber) alternating well and poorly adhered sites.

Mobility is also influenced by the interfacial structure. Well adhered inorganic inclusion barriers hinder the migration within the material, while the voids around detached interfaces provide channels for the molecular transport. The segmental mobility is mostly restricted by the interfacial interaction, but penetration of plasticizer molecules (e.g. water from moisture environment) in a surface layer can increase the local mobility. This causes a confusing increase or decrease of glass transition temperature (Tg ), depending on the character of the substrate and the polymer, when the thickness of a studied interfacial layer is decreased. These issues are of great importance because they contribute to the understanding of the mechanical, rheological, diffusion and electrical characteristics of multiphase polymers systems in which the interfacial ratio is dominant. For example a recently proposed structure with varied sites of high and low adhesion level along the surface of reinforcing fibers, as shown in Figure 4.3, is a promising concept for increasing the toughness of composites [2, 3]. The adsorption–desorption process is in close relation with the adhesion. The driving force for the adsorption is the lowering of the system’s free energy. The surplus energy of the surface of materials attracts the mobile molecules available in the surrounding medium (vapor or liquid phase). Surfaces of high surface energy (surface tension ∼500–5000 mJ/m2 , covalent. ionic, metallic bonds) are mostly covered, in normal environments, with layers of adsorbed water (already at <1% rel. humidity) and hydrocarbons. The stronglyabsorbed molecules transform these surfaces, decreasing their surface energy to that of the water’s or even lower (∼100 mJ/m2 ). Polymer or polymer-modified surfaces imply chemical control over the adsorption and desorption kinetics as the polymer identity (including side groups) and thickness influences it [4]. The hardly removable adsorbed surface layers may cause problems at processing or application (e.g. water release at processing, aggregation and decreased adhesion). The adsorbed contaminant layer hinders, for example, the adhesion in polymer-based IPs considerably [5]. Similar, but much more complex, adsorption cascade occurs at the interfaces of biomaterials when implanted into biological systems [6]. This process, referred to as the foreign-body reaction, starts with the ‘non-specific’ adsorption of proteins at the surface of inclusion by means of electrostatic and van der Waals forces [7]. The subsequent events are the adherence of disk-shaped cells to the interface, which release proteins that direct the formation of a coating fibrin matrix. Finally, huge aggregates (thrombus) of 50–200 μm size are formed around the original nuclei. Interactions at interphases involve physical and/or chemical bonds creating the adhesion between the phases and, in some cases, activation or inhibition of chain growth and degradation reactions. The number of adhesive bonds depends on the extent of contact area. For example, interpenetration of macromolecules at the IP multiplies the sites for intermolecular contact. Sum of adhesion forces and the local segmental mobility determine the micro-mechanics of multiphase systems which in turn appears in the macroscopic deformation

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of multiphase polymer systems. The chain pull-out and chain scission of connecting macromolecular chains, as well as craze formation around the interfacial zone adsorb considerable energy, when the adhesion contact is tried to be released [8]. The role of adhesive forces is expressed by the ratio of the adhesion energy and the modulus of elastic layer (Gc /E), which is small for metals and ceramics [9]. The ratio is still quite small for engineering thermoplastics, but becomes relatively large for elastomers and other highly compliant materials, including living cells [10, 11] and soft tissues. Selective adsorption of reactants and reduced entropy near the interfaces diminish the number of collisions; consequently the chemical reactions (such as cross-linking) can be inhibited by an interface. On the other hand, the presence of catalytic ions on solid surfaces activates certain reactions (such as degradation) and the restricted mobility of degradation products in the interfacial zone can influence the pathways of reactions. The formation of IP is influenced by both thermodynamic and kinetic effects. The ‘necessary’ condition for IP formation is determined by the thermodynamic driving force; while the ‘sufficient’ conditions are the kinetic factors determining the time scale that is required for IP formation. The relative roles of these factors depend on the materials and processing conditions.

4.3.1

Role of Thermodynamic Factors

To describe the structure and properties of multiphase systems requires the understanding of the structure in the vicinity of interacting surfaces, which depends on the thermodynamic interactions [12]. The general thermodynamic features are discussed considering the basic interactions occurring at the interfaces (Table 4.1). The effective contact area of phases is determined by the amount, size and form of phases, geometry (e.g. porosity) of the surfaces, thermodynamics and kinetics of wetting. The wetting is determined by the ratio between the surface tension of the phases (e.g. solid and liquid) and their interfacial tension. Acid-based interactions promote the wetting contact substantially. The wetting process cannot be described by these terms alone because the viscosity of the liquid phase, the geometry of the surface to be wetted and the external forces promoting the wetting have to be considered as well. Mixing of two polymers is accompanied by the change of free energy. Most of the thermodynamic studies of multiphase polymer systems are based on the Flory–Huggins theory which describes first of all the free energy changes of polymer solutions and melts. Complete miscibility occurs when the integration of two materials is accompanied with decrease of free energy and no kinetic factor hinders the process. The very small contribution of the pV-term to changes in enthalpy is neglected (due to the low compressibility of the components) and the entropy of mixing (e.g. conformational entropy change of folding of polymer chains and removal of segments from ordered crystalline structure), is considered negligible as well (GAB = H AB ) [13].

Table 4.1 Basic interactions at the interfaces. Interaction

Extensive value

Intensive value

Energy change

Thermal Mechanical Chemical Electrostatic

Entropy, S Volume, V Mass, m Charge, q

Temperature, T Negative pressure* , -p Chemical potential, μ Electrostatic potential ϕ

TS -pV μm ϕq

* tension

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The energy increment per units of contacting materials (e.g. segments, molecules) is 1 w = w12 − (w22 + w11 ) 2 where w11 , w22 are the internal interactions of units of each component, while w12 expresses the interaction between units of the two components occurring at the expense of the average of the former two. The total number of such contacts is x N2 zφ1 = N1 φ2 z where φ 1 and φ 2 are the volume fractions, z is the coordination number (i.e. the number of nearest neighbors), xN 2 is the total number of polymer segments in the mixture, so xN 2 z is the number of nearest-neighbor sites to all the polymer segments. The Flory–Huggins interaction parameter, chi (χ ), is defined as χ12 = zw/kT It depends on the nature of both components and it is proportional with the enthalpy change of the material thus with the enthalpy of mixing (Hm). The interaction parameter is zero for athermal mixtures, positive for endothermic mixing, negative for exothermic mixing. The higher the positive interaction parameter the larger the energy input needed for profound integration of the components. In case of two polymer phases the microphase separation is determined by the phase diagram of a multicomponent system [14]. Interdiffusion occurs in the interfacial zone when a part of the macromolecules or segments of the neighboring phases dissolve each other at the temperature of mixing. The interdiffusion can be sustained even below the critical temperature of miscibility (Tc ) (in supersaturated state) if the rate of cooling down exceeds that of the segmental mobility. Interdiffusion cannot be achieved when the critical point of miscibility (Tc ) is close to the temperature of degradation or curing. In these cases, the interdiffusion can be facilitated by the presence of graft copolymers, containing segments of both phases, at the IP [15]. The free energies at the boundary of contacting phases can be lower or higher than the surface energy of the individual phases. The first case means a thermodynamic driving force for the formation of stable IP, while in the second case the thermodynamic incompatibility has to be overcome by the energy of mixing and kinetic control [16]. The interfacial interaction between two components of a polymer blend is influenced by:

r r r

reactivity and arrangement of functional segments being present at the interfacial area; the value of the thermodynamic interaction parameter (χ ); the energy input of the mixing (separation and collision of components).

Increasing χ reduces the rate of interfacial reaction and extent of conversion, even if functional groups of appropriate type, number and arrangement are present, due to reduced solubility of the components in the formed block copolymer and to decreased interfacial volume available for the reactions. In the lack of interfacial reaction a third component can connect and separate the phases, enabling their intimate integration even if their direct contact is not beneficial thermodynamically. The changes of free energy of contacts at both sides of interlayer are more favorable in these cases than the direct contact of the

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phases. Such structure is sustained until the interlayer is stabile under the operating conditions of the system. An interlayer is stabile when it is kept together with strong internal bonds, such as:

r r

ionic or covalent bonds in solid nanoparticles (acting as connecting units between the phases); primary bonds of connecting (amphiphilic or block copolymer) molecules.

Concerning three component polymer–polymer–solid systems an attempt was made to describe the thermodynamic equilibrium in the framework of the Flory–Huggins theory [17]. The positive value of the thermodynamic interaction parameter for the ternary system (χ (A+B+S) ) corresponds to an immiscible system whereas negative to miscible one. The change in free energy by mixing two polymers and solid particle depends on the compatibility of the two polymer components. Lipatov [16] studied the free energy (G) of three components system of a solid particle S, polymers A and B: G mix = G AS + G B S − G AB In the case of immiscible polymer pairs the ternary mixture is more stable thermodynamically (their adsorption at the interface is higher than that of the miscible pair). This is the basis of the nanocompatibilization, i.e. compatibilization with the aid of nanoparticles [17, 18]. The energetic factors are mostly considered when the formation of interfacial structure is discussed, while the entropic factor is mostly neglected. This approach is reasonable, for example, in the case of thermosetting composites, where the concentration gradient of low molecular weight monomers around the inclusions (causing imperfect crosslinking) is determined by their relative affinity to the surface of inclusion. The driving force for the segregation is the lowering of the system’s free energy. In thermoplastic composites, however, the larger molecules will suffer a greater loss in conformational freedom in the vicinity of the surface; thus an entropic driving force promotes the accumulation of lower molecular weight components in the interfacial zone. Furthermore, lower molecular weight species tend to crystallize more readily than the higher molecular weight components. In such cases, the role of the entropic factor in the IP formation should be considered. The heat of mixing originates from primary and secondary interfacial bonds. The interfacial volume in which these bonding reactions occur is defined by the interfacial area per unit volume and the interfacial thickness [19]. The mixing can be, however, accompanied with other processes (crystallization, crosslinking, degradation), the kinetics of which is modified by the change of local composition and interactions at the interface. 4.3.2

Role of Kinetic Factors

Kinetic factors (diffusion, nucleation, viscoelastic and plastic deformations, biophysical response and chemical reactions) determine the time scale needed to achieve a thermodynamically favored interphase structure. Technology, processing times and application circumstances imply the relative roles of the thermodynamic and kinetic factors in the formation and response of interphases in multiphase thermoplastic and thermosetting systems [20]. The thermodynamic model of adhesion broadened by accounting the adhesion is developed not only for the thermodynamic energy of adhesion, but also for the irreversibly dissipated energy occurring in the deformed materials and at the crack tip [21, 22]. The loss of conformational freedom of macromolecules at rigid interfaces is reflected by a decrease in entropy and changes the kinetics of their response as compared to molecules far from the interfacial regions. Above Tg . the reduced segmental mobility, while below Tg , the side chain mobility can be affected. Thus

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Current Amps e-11

Current Amps e-9 -0.2

0

-0.4

-0.5

PVP

-0.6

-1.0

PVP+25%w

-0.8

-1.5

-0.8 -2.0 -1.0 -2.5

-1.2

-3.0

-1.4

-3.5

-1.6 25

50

75

100

125

150

175

T, °C

Figure 4.4 Shift of Tg peak of polyvinylpirrolidone (PVP) due to absorbed water determined by Thermally Stimulated Depolarization Current (TSC) method.

increasing the ratio of interphases, e.g. in nanocomposites, the majority of the properties are governed by entropy-controlled kinetic factors. In the case of competitive processes such as reactions leading to degradation this factor may alter the pathway of the reactions. These processes depend also upon the duration of the diffusion of components toward the surface comparing to the rate of the reactions there. Diffusion determines the environmental stability of the interphases when exposed to a moist environment. The phase borders absorb water, which can cause plasticization and swelling and consequently alters the Tg and the local stresses (Figure 4.4). Surfaces of some inclusions, such as glass fibers, are susceptible to humidity (leaching of soluble oxides) [23, 24]. Single fiber fragmentation test (SFFT) results indicated severe fiber/matrix interface degradation (increase of the critical fiber length indicating a decrease in fiber/matrix interface shear strength) for glass fiber-reinforced composites subjected to the environmental conditions. The information obtained from the single fiber tests was consistent with the significant reduction of transverse tensile and flexure strength after environmental conditioning [25]. The increasing role played by the interphase region in the performance of innovative multiphase materials continues to motivate research aimed at tailoring the IPs and calculating the final properties based on the characteristics of the IP. 4.3.3

Relationship Between Interfacial Structure and Mechanical Response

The macroscopic deformation of multicomponent polymer systems depends on micro-mechanics determined by the adhesion forces, interpenetration, interfacial mobility, and fracture performance within the boundary phase. Thick interfacial regions transform the mechanical energy to heat, through pull-out and scission of connecting macromolecular chains and formation of crazes within the zone of interfacial influence, before the adhesion contact is broken up [8]. Thus the calculations must involve several phenomena not accounted for in early models based the Kelly-Tyson equation [26]. Models considering the interface phenomena

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(interfacial layers, surface energies, etc.) appeared in the 1980s [27–29]. The interphase phenomena have been considered later by continuum mechanics models and principles of the numerical Finite Element Analysis (FEA) [30, 31]. The interfacial structures in micro- and nanocomposites and their consequences to the macroscopic properties, in accordance with the relevant modeling features, have been discussed in details by Jancar [32]. The role of the micro-scale IP in the stress transfer between the phases was proven experimentally by a single fiber fragmentation test, which exhibited strong dependence of interfacial shear strength on the local elasticity modulus [33, 34]. In polymer nanocomposites, spatial distribution of the conformation entropy within the polymer phase is of primary importance; thus the reinforcement mechanism is based on the entropy. Furthermore, the distance among the nanoparticles falls into the length scale of the radius of gyration of the polymer chains [35]. Due to the large surface area of the nanofillers, the interphase of distance-dependent gradual characteristics dominates the properties of the nanocomposites. Use of ‘reptation’ and ‘percolation’ models has been proposed for this case to describe the immobilization (instead of using the continuum mechanics models, which are valid only for microcomposites) [36]. The percolation threshold of the model is volume fraction of nanoparticles which corresponds to the case when the interconnections of the immobilized interfacial chains form a physical network within the entire volume of material [37]. The relationship between the models and the experimental results is, however, not yet well established, probably because the interfacial effects are overwhelmed by the larger-scale flaws such as voids, local variations in resin rich areas, local concentration of inclusions, residual thermal stresses, and failures [38].

4.4 Interface Modifications: Types and Aims Based on the studies for understanding of the role of interfacial interactions in multiphase systems the interface modifications for controlling these interactions (and their consequences) can be designed. The types of interface modification can be classified as shown in Figure 4.5. The capabilities of the interface modification up to now have met a wide variety of the aims of industrial sectors. Many examples of them will be shown in this chapter but the discussion of other relevant fields – such as tribology, abrasive resistance, conductivity, insulation, corrosion inhibition, irradiation screening, energy storage, flocculation, separation and membrane techniques – would need more detailed study. It is estimated that the possibilities of interface modification have not yet been utilized entirely and the field is expected to grow rapidly.

INTERFACE modification morphology

segmental mobility

compatibility

reactivity

responsiveness

transcrystalline amorphous diffusion permittivity compliance physico- mechanical insulating activating irreversible smart chem.-bio-

interdiffusion inclusion coupling

transporter converting expandable

Figure 4.5 Classification of main types of interface modification.

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4.4.1

Interlayers of Controlled Morphology

The control of the morphology in polymeric systems is one of the most convenient ways of adjusting the mechanical, optical and transport properties of the system according to the requirements. Methods for controlling the morphology are important traditionally in the technology of synthetic polymers and are the focus of current research in pharmaceutical technology. The surfaces of proper inclusions provide efficient ways to govern the morphology in their neighborhood and if these surfaces are large enough (due to the high level of dispersion of the inclusions) the control is extended to the whole material. At present, optimal selection or modification of the inclusions according to the required morphology is based on former experiences or trial and error. Better understanding of the relevant mechanisms is needed for more conscious design of interfacial morphology. Transcrystalline interlayer is widely used for nucleating the formation of preferred crystalline structures. Nucleating agents perform their action mostly through transcrystallization, which involves epitaxial growth of crystalline phase [39]. An example in polypropylene (PP) is shown in Figure 4.6. The formation of an epitaxial structure of semi-crystalline polymers, proteins, pharmaceuticals and dyes involves steps of adsorption of molecules on a surface (so that the crystal structure of the growing material can develop at or near the lowest energy state possible), formation of rodlike aggregates, their organization into crystals, and crystal growth [40]. (Wetting mechanism, desorption, interdiffusion and even chemical mechanism of ionic chain end formation by cutting of the (macro) molecules can also influence the transcrystallization processes.) Substrates having a similar nature and the same crystalline organization induce controlled crystallization in strictly defined crystallographic orientations. This kind of epitaxial orientation relationship has been explained in terms of the alignment of the chain segments of the guest-crystal-forming material along the lattice plane of the host-crystal (with an intermolecular distance of about 0.5-nm). Wittmann and Lotz have defined 15% disregistry between the matching lattice spaces of the two crystals as an upper limit of the process [41]. The intensity of the [α]-transition, related to molecular mobility associated with the presence of crystals, was found to be proportional in polypropylene to the volume fraction of nucleating inclusion [42]. Many researchers have found that the formation of the transcrystalline interlayer could improve the stress transferring from matrix to the inclusion and thus enhance the mechanical properties. However, the amorphous molecules in the transcrystalline zone immobilized by the interaction with the inclusion also play a role in composite behavior and need to be considered when interphases are tailored.

4.4.1.1

Amorphous Interlayer

The amorphous interlayers meant no interest for polymer scientists for a long time. Therefore the structural requirements for surfaces that induce interfacial amorphization are not known. Current interest for amorphous interphase exists, however, in the field of ceramics [43]. More recently, the pharmaceutical industry has started to focus on the amorphous interfacial layer of drugs. Interfacial interactions between solid substrates and drug molecules can restrict the mobility of the adsorbed molecules to an extent that hinders the arrangement required for crystallization. Amorphous drugs have a higher kinetic solubility and dissolution rate than their crystalline counterparts. The large surface area of an amorphous drug, adsorbed on the surface of a substrate, leads to an enhanced dissolution rate. Humidity-induced recrystallization is hindered by the interfacial interactions, which stabilizes the amorphous form of the interlayer. The micro-Raman technique has been used recently for detecting the interfacial amorphization and recrystallization processes [44]. A pharmaceutical composition comprising a low-solubility drug adsorbed onto a substrate of SiO2 , TiO2 , ZnO, Al2 O3 , or zeolite has recently been claimed to form a stabile amorphous drug interlayer of improved dissolution rate [45]. Amorphization (as well as depolymerization and crosslinking) can take place, owing to interfacial interaction, at blending of cellulose, chitin and chitosan with natural and synthetic polymers when

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30 μm

30 μm

Figure 4.6 Gradual development of transcrystalline layer around a talc particle in PP at 130◦ C.

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high pressure and shear deformation is applied. The results demonstrate the advantages of the method of joint action of high pressure and shear deformation compared to the conventional techniques of polysaccharides mixtures production [46]. Extrusion was reported to be an efficient way of joint action of high pressure and shear deformation for controlling the morphology of interlayer in pharmaceutical compositions [47]. 4.4.2

Interlayers of Modified Segmental Mobility

Interaction of phases modifies the segmental mobility in the IP, which is reflected in macroscopic properties such as diffusion coefficient, permittivity and compliance. Increasing the segmental mobility makes these values higher, while reduced when the mobility is restricted. The size and structure of the interlayer as well as the strength of phase interactions have to be taken into consideration when these properties are to be controlled by the addition of inclusions to the polymer system. Beyond the characteristics of the polymer matrix the surface composition and geometry of the inclusion, the processing parameters and the interfacial additives determine the interface-relating properties including the diffusion coefficient, permittivity and compliance. Diffusion coefficient (D) expresses the value of diffusion-ability of gases or liquids through solid substances. Control of the segmental mobility determines this transport through polymers. Modification of interfacial interactions affects the degree of crystallinity and hinders or enhances the segmental mobility in the amorphous phase resulting in increased or reduced diffusion coefficient. For example, enhanced water diffusion occurs in the case of weak interaction at the interfaces. Nanoparticles of barrier, nucleating and/or immobilizing effect are introduced for limiting the transport through packaging materials [48]. A remarkable slowing of the diffusion has been observed as a result of restricted segmental mobility, even at distances greater than 10Rg from the solid surface [49, 50]. The lateral diffusion of polymers in thin layers shows linear reduction in D as a function of the layer thickness [51]. Permittivity (ε) is a measure of the ability of a material to be polarized by an electric field. The dielectric constant (relative permittivity, k) of a material is the ratio of its permittivity to the permittivity of vacuum. The dielectric properties of multiphase materials play an important role in achieving the desired performance of integrated circuits (IC, e.g. semiconductor fabrication and packaging). Interphases of multiphase systems are serving for fine tuning of the dielectric properties of the matrix polymers and additives. Molecular structures in these regions mediate the charge motion affecting the charge injection and trapping in semiconductors [52]. Thermally stimulated depolarization current measurements have shown that deeper energy states can be engineered at the oxide and polymer interface by incorporating interfacial polar charge withdrawing functional groups [53]. Low permittivity (low ability to polarize the segments of the polymer) is required for good insulators in dense multi-layered ICs, while high permittivity is preferred for capacitors and memory cells. The frequency dependence of the permittivity of polymers reflects the fact that their polarization does not respond instantaneously to an applied field. Dielectric spectroscopy methods, such as measurement of thermally-stimulated depolarization currents, utilize this feature to characterize the segmental mobility and interfacial structures in materials [54]. Voids at the interfaces act as local capacitors (Maxwell-Wagner effect), while the presence of polar molecules/segments (e.g. water) in the interphase modulates the response of material to the electric field [55, 56]. The role of IP increases substantially in the case of nanocomposites. If there is anisotropy in the shape of the phases, the difference in the dielectric constants of the domains will force an orientation of the domains in the direction of the field lines. Although the field necessary to overcome the thermal randomization increases considerably as the particle size decreases, alignment of anisotropic nanoparticles (e.g. montmorillonite, carbon nanotube) by electrical field is feasible [57–60]. Epoxy nanocomposites can be adjusted this way for the needs of advanced sensors and microelectronic devices [61]. Ultrahigh-energy density capacitors with microsecond discharge and ultradurable electrical insulator materials can be designed utilizing different types of interfacial structures [62, 63]. Nanoporous materials of dielectric constant lower than 2.0 are promising for insulating extra small wires and transistors

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but the role of the mobility of interfacial segments in the formation and stability of nanopores has to be understood. In addition to the uniquely tailored dielectric properties the controlled thermal expansion, heat conductivity, fire retardancy of nanocomposites facilitates their electronic applications too [64, 65]. Compliance (deformability) is the reciprocal value of the elasticity modulus. Segmental mobility of the macromolecules in the IP, modified by the interfacial interactions, strongly influences the compliance. Elastic properties of IP are controlled by the adhesion and the conformation entropy of the chains forming the layer. Mobility restricted by the interactions limits the compliance, while larger mobility increases it (decreases the modulus). In thermosetting composites, depending on the inhibiting or activating effect of the inclusion on the crosslinking reaction, the network density in the IP (and consequently the elasticity modulus) can be decreased or increased [66]. In thermoplastic nanocomposites all the partially entangled macromolecules at the surface are in contact with each other [67]. The extended immobilization at the enhanced specific surface area causes a steep decrease of compliance (increase stiffness). Stiff interphases provide efficient stress transfer and hinder the diffusion of water; however, support brittle failure thus limits the damage tolerance. Too rigid interlayers do not release the local stresses when external load or thermal shrinkage difference occurs. These local stresses accumulate in the interfacial region and cause debonding when the level of adhesion is reached. Such interlayers can be considered, mechanically incompatible’ with the phases. In contrast, if controlled segmental mobility in the vicinity of phase border is preserved it distributes the local stresses in larger volumes (by crazing and plastic deformation), resulting in increased fracture toughness and maintained Young modulus [68]. Tough interphases slightly reduce the effectiveness of stress transfer but provide significant enhancement of damage tolerance and reduce the sensitivity of anisotropic composite structures to the direction of the external loading. Interlayer of controlled segmental mobility is achieved mostly by introducing suitable compatibilizer IP between the phases. 4.4.3

Interlayers for Improving the Compatibility of the Phases

Interface modification, aiming at improved harmonization (compatibilization) of phases within multiphase systems, includes not only the wide area of polymer composites and blends but also the strongly developing field of biocompatibility/bioadhesion. Interfaces, having low interaction among the adjacent phases, do not bear the stresses originating from external loading or differences of shrinkage. As a result of these effects, the stress concentration at the interface leads to separation of the phases from each other and void formation around the inclusions. The mechanical and diffusion barrier properties of such IPs are obviously poor. A general rule for IPs of disperse systems was defined 60 years ago by the Ostwald-Buz´agh continuity principle [69], which postulates optimal IP as a case when the transition between the adjacent phases is ‘harmonic’ (regular) rather than sharp. A hierarchic multilayer IP structure was proposed on this basis for polymer composites [70, 71]. One of the ways of forming these multilayer interfacial structures is the arrangement of surfactants and elastomers around filler particles [72]. The IP around the inclusion, according to a model describing this structure, consists of a thin layer of oriented surfactant molecules and a relatively thick elastomer layer. It was found that such IP promotes the homogeneous distribution of the components and the maintenance of interaction with the polymer matrix and thus is suitable for balancing the stiffness and toughness of polymer composites [73]. Orientation of macromolecules at the interface in one direction (owing to the local shear strain) improves the strength of composites [74]. 4.4.3.1

Physico-chemical and Biocompatibilization

The chemical and biological differences between the neighboring phases have to be bridged if longer-term interaction is required. Even if thermodynamic compatibility (leading to molecular dissolution) can be hardly

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achieved the modifications ensure sustained ‘kinetic compatibility’ (restricted segregation-mobility), which is satisfactory in most cases. Natural biocomposites provide a challenge for designing interfacial structures. Interphases at different length scales in hard tissues like bones, teeth and shells are hierarchical, adaptive and self-repairable [75–78], resulting in mechanically stiff and tough natural composites Artificial biomaterials must also perform harmonic transition at the biomaterial–tissue interface but the main aim of the coverage of biomaterials is to minimize the immune reactions [79]. For this purpose, biomaterials have to be surrounded with soft, hydrophilic macromolecular IP, which inhibits the spontaneous adsorption of non-specific proteins (fouling) [80]. Methods developed for forming non-fouling soft, hydrophilic surfaces use generally the following substances: liposomes immobilized through functionalized coupling molecules (polyethyleneglycol, biotin, avidin), amphiphile polymers (containing ethylene oxide units of min. 500–1000 molecular mass), methacrylatepolyethyleneglycol (PEG) oligomers, hydrogel of functionalized polylactic acid and PEG, metacrylatepolymer containing phosphatidyl choline, polysulfone, poly (N-vinyl pyrrolidon), polyacrylamide, etc. Such modifications can be initiated also by high energy physical (such as UV, ion beam, plasma, physical vapor deposition) treatments [81]. As an example, plasma deposition of tetraethyleneglycol dimethylether shown to have extremely low protein pickup due to the formed crosslinked PEG-like structure [82]. Polymer composites and blends require IPs of controlled adhesion, which is performed through the following ways:

r r r

interdiffusion of the macromolecules of the attached phases; inclusion of amphiphilic molecules, block copolymers or nanoparticles into the IP; and reactive coupling with functionalized polymers, coupling agents, reactive surfactants.

Interdiffusion at the interfaces can be promoted by thermal energy, mechanical energy and chemical modification. The thermal energy increases the solubility of the phases in each other in the cases of upper critical solution temperature (UCST) systems, while the mechanical energy increases their contact surface through reducing the particle size. The use of cosolvent facilitates the interaction of the phases However, the complete drying of solvent from polymer would take a very long time, and sometimes is even impossible because of the interaction between solvent and polymer. Therefore a compressed gas, such as CO2 , for promoting the interdiffusion is more favorable [83]. An interpenetrating network (IPN) structure is formed when forced interdiffusion of thermodynamically incompatible phases is followed by separation limited by (chemical or physical) crosslinks, which prevent the full phase separation. IPNs are in the state of quasiequilibrium but the true equilibrium, due to kinetic reasons, can be never reached. Inclusion of amphiphilic molecules, block copolymers or nanoparticles into the IP region can promote the interaction of the phases. Their role is to overcome the physical–chemical (polarity, charge, etc.) differences at the IPs of polymer blends and composites by acting as wetting agent (surfactant), performing interdiffusion into both phases (block copolymer) or providing enlarged adhesion sites for both phases (nanoparticles). The main role of ionic and non-ionic surfactants in (nano)composites (e.g. in oganoclay-containing systems), is to promote the wetting and the homogeneity of components in the (nano)dispersion [84]. Block copolymers acting as ‘connector molecules’ or at interfaces of polymer blends or layered structures serve as sites for interdiffusion [85]. The inclusion of the nanoparticles in the interfacial region initiated the concept of nanocompatibilization. One of the approaches of nanocompatibilization is the formation of scarf joint bonding through the arrangement of carbon nanotubes (CNTs) along the interface of the composite joints. It was found that an optimal amount of carbon nanotubes could increase the fracture toughness of the composite joint interface significantly by modifying the fracture mechanism [86]. Another method for nanocompatibilization is the introduction of nanofillers into an immiscible polymer system leading to the formation of a thermodynamically stable system. Organoclays and other nanoparticles have been proven to

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Figure 4.7 Scheme of compatibilization of immiscible polymer phases with nanoclay.

improve the performance of incompatible polymer blends. Polymer chains of both polymer phases can be absorbed to the clay surface to form more compatible nanocomposite [18]. Organoclays of high aspect ratio (500–1000:1) and large surface area (700–800 m2 /g) create sites for interaction with polymer chains. Beyond the thermodynamic factors the increased viscosity and local shear forces in presence of nanofillers contribute to the improvement of the miscibility of incompatible polymer phases. The polymers may have different interactions with organoclay. For instance, in polystyrene/ poly(methyl)methacrylate (PS/PMMA) blend the organoclay is preferred by PMMA, while in polycarbonate/styrene acrylonitrile copolymer (PC/SAN) blend the organoclay has no preference. The large surface-tovolume ratio is efficient in reducing the interfacial tension between the phases in both cases. The (dark) clay platelets are observed to bend in order to conform to the contours of the interface as shown schematically in Figure 4.7. Equilibrium morphology is determined by the reduction in interfacial energy and the increase in bending energy [18]. Application of this compatibilization concept is obviously advantageous in the case of recycling of mixed polymer wastes. Less obvious but also useful is the concept in the case of fire retarded polymers. Fire retardancy (FR) of polymer systems containing polypropylene (PP) and phosphorylated epoxy resin (PEP), was reported to increase when montmorillonite nanoclay was introduced. Without nanoclay the distribution of epoxy in PP matrix was inhomogenous, while homogeneous dispersion could be achieved after attachment of PEP to clay nanoparticles as interlayer. The improved FR performance is ascribed to homogeneous char forming action of PEP interlayer [87]. Combination of recycling and fire retardancy concepts in the form of fire retarded recycled nanocomposite allowed upgrading recycling (upcycling) of polymer wastes [88]. Hierarchic IPs, containing inclusions of more than one type are, in many cases, indispensable for solving the interfacial problems of multicomponent polymer systems [89]. The described strategy of physical compatibilization of phases in multicomponent systems, by including non-reactive substances into the interfacial layer, is limited by the strength of secondary bonds. The subsequent reactive compatibilization strategy involves primary bonds for increasing the strength of interaction. Reactive coupling is the most effective way for enhancing the adhesion between the phases improving thus the stiffness/strength and other (such as barrier and stability) characteristics considerably. Extensive literature reviews are available on the reactive compatibilization of multiphase polymer systems [90–92]. In the course of such IP modification, reactively bonded substances must replace the spontaneously adsorbed weak boundary layers (that induce failures especially under combined conditions of high stress and humidity). In situ reactive compatibilization can take place either by reacting two types of functionalized macromolecules during the melt processing or by using bifunctional molecules (e.g. coupling agents) for reacting with both

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phases. The functional groups should be selected so that the interfacial reaction occurs within the time frame of processing time (typically < 5 min). The efficacy of in-line IP modification depends on the following factors:

r r r r r r

miscibility of the active molecules with the phases; mobility of reaction partners (so that they reach the interface rapidly); reactivity of functional groups (as determined on small molecular models); structure of reacting molecules (graft is slower than end-functional reaction) [92]; stability of interchain bonding (against the thermal and shear stresses of processing); compounding parameters (∼1000 times higher reaction rate under flow than in static bilayer) [92].

Under higher shear of compounding the surface area is increased by breaking down the dispersed phase, which in turn increases the probability for functional groups to meet and react with each other (the reaction reduces further the phase size and forms barrier action against coalescence). It was found that the reaction rate is accelerated by decreasing thermodynamic interaction parameter χ [93]. Functional groups promoting compatibilization (without by-product) include maleic anhydride, NH2 , OH, SH, COOH, CHO, and epoxy groups (Figure 4.8). The most commonly-used reactive pair is primary aliphatic amine with cyclic anhydride [90, 91]. The most rapid reaction occurs between amine and aromatic isocyanate but the formed urea bond is sensitive to water and high temperature (which in turn is advantageous for thermally reversible crosslinking). Compared to macromolecules functionalized with reactive groups, such as maleated polypropylene (MAgPP), the relatively small coupling agent molecules (reacting both sides of neighboring phases) are more mobile. Some examples of the characteristic silane coupling agents used to improve the interfacial adhesion in multiphase systems are shown in Figure 4.9. In some cases, higher content of functional species are added in order to ensure the required performance. However, the reactive groups which are not converted during the reactive compatibilization can further react during subsequent processing (re-extrusion or injection molding), which might result in deleterious consequences. This clearly means that the kinetics and yield aspects of the interfacial reaction in reactive compatibilization are of the utmost importance. Many coupling agents are immiscible with the polymers and form their own phase. Additionally, partial crosslinking of coupling agent (e.g. hydrosilane) may occur under mixing conditions [94]. The parallel coupling and crosslinking processes are affected by substrate, external energy source and by the atmosphere (e.g. oxygen, humidity) being present. Plasma-assisted methods combined with silane coupling agents produce polysiloxane interphase of higher network density (and higher modulus) than the solution deposition techniques. Enhanced fire retardancy could also be achieved by plasma treatment of polysiloxane layer [95, 96]. The energy of an electron beam, applied in course of melt blending, may lead to reactive compatibilization without any additive (Figure 4.10) [97]. It is not easy for the interface modifiers to reach the right interface during the compounding process. Surfactants, due to their surface-active character meet this requirement but for coupling agents it can be safely ensured only by a separate surface treatment step. Reactive surfactants (RS) have been developed for combining the benefits of surfactants and coupling agents [98, 99]. These interfacial additives, considered as targeted reactive interphase modifiers, tend to find the interfaces, due to their amphiphilic character, within a short residence time and bond the phases preferably. Such additives containing carboxylic anhydride groups in the polar region and reactive double bonds in a non-polar hydrocarbon chain have been synthesized by means of Diels-Alder reaction or esterification [100]. Dienophile compounds, e.g. maleic anhydride was used as reaction partners to the Diels-Alder reaction. Use of RS additives resulted in improved mechanical properties in polymer blends, while in pigmented PE it contributed to higher photostability [101].

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Without residual by-product O OH

HO R′

H2N R′

O R′

O R

R

O O H 2N R′ or: O C N R′

O

HO

HOOC R′

epoxide

N R′ anhydride

O

R

O

R

H N R′

HO

R′

O

R

R

O

O

N H

R HOOC R′

O R

O

R′

N

H N

O

H2N R′

O

O R

R′

HOOC R′

isocyanate

2 HOOC

R″

O

R'

R′

R C C R′ H H

O

R″ ;

R′′′

R N C H

O R′ O

With by-product O R

O

H 2 N R′ R

OH

R Si OR′ OR′

silanol

HO M

N

R′

+

H2O

H

carboxylic acid

n OR′

R

and H2O O

surface with free hydroxyl groups

R″ O

N C H

R″ O

O

double bond

HO R′

R N C H

H N R′

R N C H

- CO2

carbodiimide

- CO2

R N C H C O

C N R′ H R″

C N R′

N O

H2N R′

R N

R N

O O

O

R N

HN

N H

R

H2N R″

R′

O

O

oxazoline

R′

R

Si O

Si O

M

M

+ 2n HO

R′

n

Figure 4.8 Reactive groups commonly encountered in reactive compatibilization

H R′′′ C C H R″ R′ R

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O

H3C

CH3 O

O Si C CH2 H O H3C

O O

H3C

C H2

H2 Si C

H3C

H2 C

N H

H2 C

O

H3C

H2 C Cl C H2

C H2

NH2

O

O

C H2

O H3C

H2 C

O

H3C O

O

H3C O

SH

H2 Si C

O

H2 C O C H2

H3C methacrylate

H2 Si C

O

O

H2 H2 H C O C C O C C H2 H2

CH3

H2 O Si C

CH3

O

C H2

H2C

CH2

NH2

epoxy

H3C

O

CH3

H3C

thiole

CH3

C H2

H2 C

primary amine

H3C diamine

CH3

H2 Si C

O

CH3

H2 Si C

O

H3C

chloropropyl

CH3

H2 O Si C

CH3

O

vinyl

H3C

O

H3C

HC

H2 C

H2 C NH2 CH3

+ Cl

-

CH CH

HC C H

cationic styrile

CH2

Figure 4.9 Examples of silane coupling agents.

Thickness of the modified interlayer falls in the range of 1.5 to 500 nm depending on the treatment technique. Extensive interpenetration between the coupling agent and the matrix monomer in thermosetting matrices leads to a thicker interlayer, which is not characteristic to the IP in thermoplastic polymers. In the case of epoxy composites containing E-glass fibers treated with γ -aminopropyl-triethoxysilane (γ -APS) the bond strength increased 1.7 times and toughness 1.9 times higher compared to unsized samples [102]). Such interlayers are not monolayers but thin IP of network polymers crosslinked after deposition by the condensation reaction of hydrolyzed Si-OR groups of the coupling agent. Gradual decrease of the stiffness of the deposited IP was detected as a function of the layer thickness reaching a bulk value for layers thicker than 105 nm [32, 103]. Reviews on the effect of surface-grafted molecular brushes on the adhesion performance of polymer composite interfaces have been published recently [15, 104]. These structures, after their synthesis, can be used in reactive processing technologies and form in most cases high level of adhesion and gradient IP structure (in which the modulus of elasticity changes across the interfacial layer gradually) [105]. Instead of the gradual change of the interlayer the more complex core-shell structure is favored in a number of multiphase systems for optimizing the toughness/stiffness balance of composites’ mechanical properties

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accelerating voltage vacuum catode Wehlnet unit anode focusing adjusting die

screw

Figure 4.10

In-line electron beam treatment during extrusion.

(see mechanical compatibilization). The synthesis of covalently bonded core-shell layers on nanoparticles was successful using the ATRP (atom transfer radical polymerization) method [106]. A variety of inorganic substrates can be functionalized this way with well-defined block copolymers possessing both glassy and rubbery segments. For example, brushes polymerized directly on the surface of nanoparticles resulted in a hard colloidal core surrounded be a soft inner shell and glassy outer shell [107, 108]. According to this approach block copolymers containing segments of different polarity can be attached on functionalized surfaces that respond to the changes of environment by structural rearrangement shown in Figure 4.11. The difficulty of achieving high grafting densities, due to steric hindrance of adjacent grafts, can be overcome by stretching the substrate before growing brushes under ATRP conditions. The strain can then be released, allowing the substrate to return to its former size, producing densely grafted polymer brushes [109]. RAFT (Reversible addition–fragmentation chain transfer polymerization) is similarly or even more versatile than ATRP for the production of complex architectures. It uses thiocarbonylthio compounds in order to mediate the polymerization via a reversible chain-transfer process. This technique was applied for synthesis of brushes displaying reversible surface properties upon treatment with different solvents [110]. The ‘click chemistry’ principle is also applicable for interface modification. A layer of 20 weight % around cellulose ester fibers,

polar apolar R R x x x x x x x

x

Figure 4.11 Formation of smart interlayer of apolar (o) and polar () segments and its arrangement in polar and apolar environment.

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bearing multiple C–C-terminated hairs, was formed by polycaprolactone-diol (PCL). The PCL was converted to azido-derivative and the cellulose esters were reacted with azido-PCL grafts [111]. 4.4.3.2

Mechanical Compatibilization

Macromolecular IP-s are present in most of the widely-used multiphase polymer systems. Polymer-layered metals, reinforced polymer systems, polymer blends and various kinds of biomaterial implants can be found among these types. Mechanical properties of filled and fiber reinforced composites depend on the micromechanical properties of the interphase zone, which is affected by the adhesion, thickness and structure of the IP. Harmonic transition between the phases requires the differences of their elasticity to be bridged without which the stress concentration at the phase border exceeds the adhesion forces. The reinforcing effect of polymer composites of mechanically incompatible interfaces is limited by the resulting dewetting of fillers, pull out of fibers or delamination of layers. Introduction of an intermediate thick macromolecular IP is needed in order to avoid stress-induced debonding and void formation. Macromolecular IP-s around fillers and fibers can be formed by off-line technologies, for example aerosol and solvent-based procedures used for forming well-adhering coatings of polydivinylbenzene on silica beads [112]. Much more simple, however, is the ‘in situ’ process that has been proposed for forming relatively thick (100–500 nm), soft elastomer interlayers during a single step compounding. The principles affecting the arrangement of the elastomer phase in the composites, i.e. its share between the interphase and disperse phase, was published in the 1980s [72, 113, 114]. It was found that the concentration, polarity, viscosity and melting regime of the IP-forming elastomer determine the thickness of the IP containing a strongly immobilized stiff ‘inner’ layer and an ‘outer’ layer of high toughness. Thus the structure–property relationship of the composite can be governed by the thickness and structure of the IP [115]. The elastomer in the boundary layer and the coated filler particles participate coordinately in the toughening mechanism of such multiphase systems [115]. The combination of stiff interphases (providing efficient stress transfer) and tough interphases (decreasing the damage sensitivity and strength anisotropy) is advantageous. It is provided by hybrid fiber reinforced composites (FRCs) in which the reinforcing fibers are coated with both stiff and tough interphases in order to tailor the performance and reliability of the final FRC part [32, 116].

4.5

Interlayers of Modified Reactivity

Reactions of substances in the interfacial region are not restricted to the discussed coupling reactions and the pathway of these reactions can be influenced by the chemicals being present at the phase borders. These interfacial reactions are required to suppress or promoted (depending on their type) by modification of the surfaces of the components. Scission of the instable groups of molecules adsorbed at solid surfaces can be promoted by altered electron density. A part of the additives having catalytic ions on their surfaces initiates degradation at elevated temperature. Surface of montmorillonite and attapulgit contains, for example, traces of catalytic iron ions, while colloidal silicone can act as Lewis acid catalyst [117]. The moisture, adsorbed by the IP, reduces the barrier effect of hydrophilic polymer and decreases the stability by initiating hydrolysis or acting as solvent in an interfacial reaction [117]. Main reactions leading to degradation are hydrolysis, oxidation, isomerization, polymerization and photolysis [117]. Photolysis can be suppressed by introduction of nano-TiO2 into the interfacial region but coating of its surface reduces its capability for modifying the photochemistry [117]. The stability of amorphous interfacial region has recently been of special interest. The amorphous phase is more sensitive to chemical, thermal, UV, and microbiological degradation than the crystalline one. This obvious statement, which refers to the case when the phases are examined by themselves, is the source of

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misunderstanding in many cases. The amorphous phase, if it is insulated from the initiators of chemical reactions (such as O2 , H2 O, UV etc.) by the appropriate interlayer, can be even more stabile than the uncovered crystalline one. The amorphized (and other sensitive) components therefore have to be protected by an insulating interlayer. Insulating interlayers are designed to separate components from each other and thus prevent undesirable interaction. Examples of industrial importance are the coatings applied to avoid fouling or degradation of phases [118]. Insulating interlayers are applicable to overcome the chemical and thermal sensitivity of fire retardants [119]. Ammonium polyphosphate (APP) fire retardant additive, for instance, has to be protected against hydrolitic degradation at higher temperatures, for which purpose several insulating polymer layers have been developed [120, 121]. In the case of a shear stress resistant silicone elastomer the protecting effect (preserved under processing conditions) was confirmed by the decrease of the conductivity of water used for extracting the FR product (from 200 μS of the reference sample to 50 μS of protected APP) [119]. Activating interlayers are modified in order to enhance catalytic activity of interlayers. Enzymes of controlled catalytic action have been immobilized on solid surfaces by means of sol-gel reaction [122]. Activating interlayers can be utilized also for controlling the mechanism of fire retardants. For example melamine polyphosphate-polyol system, which is stable at the degradation temperature of PP, has been activated for application in polyolefins [123]. An activating effect is ascribed also to the char forming polymers in intumescent nanocomposites, according to the blending approach [124].

4.6 Responsive Interphases The IPs discussed up to now possessed static function. However, the role that an IP has to play can change according to the change of circumstances. Responsive IPs are designed for accommodating to such changes. No classification is made, in general, within the responsive IPs although the irreversible or reversible character of the response is of a great difference. Therefore, in order to distinguish these cases from each other, we call the irreversibly and reversibly responding IPs non-reversibly adaptive and smart (reversibly adaptive) respectively. 4.6.1

Non-reversibly Adaptive Interphases

In non-reversibly adaptive IP a transformation, induced by an external signal, occurs at single occasion [125]. Heat-, light- and pH-induced transformation may facilitate stabilization (mechanical, photostabilization and fire retardancy), while interlayers, being responsive to the change of the biological environment, can be utilized in advanced (bio)composites. Adaptive changes involve induced transporting, expanding or transforming actions. Transporter IP act as a carrier layer, which is activated by a signal of the changing environment and adsorption or desorption processes are involved in the action. Non-reversible adaptive IP promotes this transportation in one direction. Advanced examples of the carrier coatings can be seen in nature, where hormone-based chemical signals initiate the attachment of carrier molecules on substrates, which facilitates their transport through membrane barriers. In this case, one-way transport is the result of the changed surrounding of the substrate. Controlled release from drug delivery systems can be accomplished using thermally and chemically adaptive, enzymatically erodible IP-s. Enhanced temperature and acidosis in the neighborhood of pathologic processes allows nanoparticles covered by heat- and pH-sensitive polymers (such as copolymer of poly (N-isopropylacrylamide) and polylactic acid) to release drugs at that areas. Special enzymatic environment degrade the delivery system at the target site if environment-adaptive IP supports the controlled release of

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Compatibilizer layer

Char catalyzing layer

Charring polymer layer T-responsive coupling Polymer-matrix

(a)

(b)

(c)

Figure 4.12 Thermally responsive fire retardant interlayer with chemical and mechanical compatibilizer layers at ambient temperature (a), initiated charring at the critical temperature (b), the whole interlayer transformed to protective char-foam (c).

drugs. The surface layer degrades at the target area gradually in presence of these enzymes [126]. In vivo environment-initiated NO release from polymer coating of implants helps to avoid the negative consequences of cardiovascular surgery [127]. Similar concepts can be utilized for stabilization purposes. Stabilization against degradation caused by UV light, fire, or bio-chemical impacts, requires surface protection. The stabilizers are generally homogeneously dispersed in the material but when the impact increases these substances should accumulate at the surface in order to enhance the protection function. Enhanced photostability of polymers can be achieved by modification of the surface layer through heat- and/or light-induced migration of stabilizer molecules. As the process contributes to the removal of the stabilizers from the system the reservoir capacity of the bulk means the limit of such protection concept [128]. Fire retardancy of polypropylene has been improved recently by forming non-reversibly adaptive IP around flame retardant particles [129]. The main role of IP in fire-retarded systems at normal environment, similar to other multiphase systems, is to enhance the compatibility, facilitate the processability and the mechanical properties [130]. At elevated temperature caused by fire attack, however, the IP-s around the particles should promote the development of heat insulating protective char-foam and deliver the particles to the surface where a fire-protective barrier IP is getting formed. Thermally induced detachment of a compatibilizing unit from the surface of fire retardant makes the underlying incompatible polymer uncovered which, acting as a carrier layer, delivers the flame retardants to the surface [131]. Above the critical temperature the whole interlayer expands, forming fire protective foam. Thus the surface phenomena, and especially the adaptive interface, can play an important part in the fire retardancy performance of polymeric systems. Converting IP plays an adaptive role by performing chemical or biological conversion when it is induced by thermal, chemical or UV light signal. Temperature-induced ceramization, for example, can transform a boron containing polysiloxane surface layer fairly resistant against the heat and oxidative effect of fire [95, 132–134]. Silicone-modified clay nanoparticles, utilizing similar mechanisms, contribute to the fire retardancy in polymer nanocomposites [119]. Crosslinking, induced by external influences contributes to the formation of durable protecting polymer IP. It is well known, for example, in the technology of UV-curable hurricane-resistant glass laminates. Biomimetic modification of many advanced biocomposite requires welldesigned adaptive IP. Change of the structure of IP-s according to biological environment is utilized in

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bone repair and regeneration, wound dressing, artificial skin, etc. In order to elicit the least foreign-body reactions and good tissue-bonding capability physiologically active surfaces of selective adsorbability are required [135]. In nature, specific proteins (of fixed conformations and orientations) act as signaling agents, while a non-physiologic proteinaceous layer is considered as a foreign invader that must be walled off [136]. Therefore, biomimetic modification of IP-s must control the type of conformation and orientation of proteins and thus form chemical topographic and visco-elastic patterns on surfaces. A promising approach is the attachment of peptides containing domains of proteins of extracellular matrix (ECM, natural environment of cells) to the surface of biomaterials. Hierarchic IP-s including these active peptides can induce tissue formation conforming to the cell type seeded on the material. Biomimetic surface engineering for bone regeneration can affect bone cell adhesion and preferentially induce mineralization on the modified surfaces. In the bone regeneration process the IP is required to transfer the mechanical charge to the bone gradually Adaptive IP between bones and implant promotes this process by its own gradual conversion. The presence of hydroxyapatite on the modified surface appears to have the ability to induce bone formation. The design of IP must learn from the biominerals, which are natural examples of biocomposites having unique interlayers between the mineral and organic units. Molluscs, for example, synthesize nacre of 95% calcium carbonate and 5% protein, which is mechanically tough due to optimal IP between the phases [137]. Controlled release of the enclosed drug from polymer implant or drug delivery systems can be accomplished using thermally and/or biochemically adaptive (for instance, biodegradable) IP-s. Enhanced temperature and acidosis in the neighborhood of malignant tumors induce transformation of barrier layers allowing drug release at those areas. Other types of surface layer degrade gradually in presence of certain enzymes occurring just at the target area where the encapsulated drug should act [138]. In vivo environment initiates NO release from those polymer layers that are applied as a partial replacement of endothelial cells helping to avoid the negative consequences of cardiovascular surgery [139]. Expandable interlayers perform their role by causing volume change in the material. Irreversible expansion is initiated by chemical reaction, i.e. curing or decomposition. Interface-initiated changes of thermal–mechanical properties of epoxy nanocomposites could be achieved by interlayer expansion. Differences in the activation energies of interlayer expansion and of curing influence the final nanostructure and thus the thermal–mechanical properties. The stages of the formation of interphase layer, composed by the epoxy matrix plasticized with surfactant chains, around nanoparticles (i.e. the interlayer expansion/exfoliation mechanism) determines the ultimate properties of the material [140]. The presence of montmorillonite (MMT) nanoparticles influences the foaming process of modified epoxy resin around nanofillers advantageously [141]. This effect can be utilized for enhancing the efficiency of fire-retarded systems. The particles, expanded as a result of heat and catalytic effects, may reach each other leading to percolation of the layer coated particles. Compared with the conventional intumescent fire-retarded systems, in which continuous char layer is formed only after the decomposition of a considerable amount of matrix polymer, the expanding particles of a ‘percolating intumescent system’ reach each other in an earlier stage of the process providing rapid protection to the underlying polymer phase [142]. Parameters influencing the percolation are the interparticle distance, volume increase and the rate of action [143]. Expandable interlayers provide also a way for rapid delivery of active substances to the surface [129, 144]. A material utilizing this mechanism contains interlayers of relatively low decomposition temperature intercalated between layers of nanoparticle. At the early stage of fire action the interlayer starts to decompose and the gaseous degradation products rapidly drive the separated nanolayers to the surface. 4.6.2

Smart Reversibly Adaptive Interphases

Responsive interlayers that ‘communicate’ with their environment through a reversible recognition-response mechanism are considered smart IP. The applicability of such interfacial structure requires rapid and

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considerable shifts of the surface properties (such as expansion and collapse) induced by slight external changes. Continuous development in this area is initiated by nature, which created various types of functional, controllable interfaces during the course of evolution. Ultimate ‘smart’ materials are, for instance, the receptors and trans-membrane proteins that adjust their function in response to environmental changes or biochemical influence. A smart IP generally consists of stimuli-responsive polymer (SRP), the structure and behavior of which are influenced by external signals. The generated change can initiate feedback to the environment. Both the kind and intensity of such ‘communication’ are determined by the characteristics of the modified interface [145]. The governing external signal includes T, pH, redox-potential, light, magnetic, electrical changes or even antigen-antibody interaction. A variety of approaches have been developed for forming controllable surfaces; many of these address medical strategies such as bioseparation, drug delivery, biosensor design, tissue engineering, protein folding, and microfluidics [146]. Surfaces of molecular recognition capability are formed by coupling to enzymes or other specific biological receptors [147]. Recent results demonstrate the applicability of smart biopolymer surfaces for carrying out reversible micropatterning of proteins [148]. A responsive coating layer can control the activity of the underlying material; thus its dissolution rate, catalyst action, sorption activity, ion-exchange, geometrical dimension, and conductivity can be influenced externally. Responsive hydrogel layers around gold nanoparticles have been used for controlled drug delivery and also for producing microlens of aperture and focal length controllable in seconds with IR light, which regulates the curvature of the interface by inducing volumetric change [149, 150] Smart polymeric coating layers consisting of either reversibly soluble-insoluble (SIS) polymers (such as poly(N-isopropylacrylamide), crosslinked hydrogels or their combinations have been shown to respond to a variety of stimuli. The pH-responsive coatings could be subjected to ionization of a free carboxyl group (e.g. in methacrylates), the thermosensitive coatings would deswell as the temperature is raised beyond lower critical solution temperature (LCST) providing a basis for developing smart membranes [151]. The formation of pH-sensitive membranes is an important area where smart IP can be applied. The surfaces of pores of membranes are grafted, according to a recent approach, with polyacrylic acid. These chains undergo fast and radical conformational changes with the pH controlling the way they transport through the membrane [152]. A similar method was proposed recently using layered silicate (Na-MMT) surrounded by polyorganosiloxane (Sil) derivative and the system was applied as a coating layer of membranes covering the pore surfaces. Being the change between the forms of Na-MMT and H-MMT reversible the clay renders such a membrane pH responsive. At higher pH the distance between the layers increases, the coating layer becomes swollen and the pore closes. At lower pH, however, the layers are attached to each other, the coating layer is not swollen and the pores are open. At higher temperatures (above ∼120◦ C) the pores will be opened independently of the pH [125]. An insulin-releasing hydrogel layer around a container utilizes the activity of glucose oxydase enzyme to change the pH at enhanced glucose concentration, which in turn opens channels by shrinkage of the pH sensitive hydrogel (Figure 4.13) [153]. Similar controllable IP can be formed on the surface of particles. As an example, by postgrafting of hydrophilic polymer to grafted hydrophobic chain on carbon black, amphiphilic carbon black was obtained, the dispersability in solvent and the surface wettability of which were readily controlled by pH and temperature [154]. Polymer surfaces that can alter their wetting characteristics according to the humidity of the environment have been formed by introducing a hydrophilic group to the end of surface-active blocks, which are present or hidden at the surface depending on the moisture content of the air [155]. Very active research is expected in the near future to investigate the application of smart IPs in the area of bioelectronics and artificial photosynthesis. The challenge for both areas is to form chemically and topographically patterned surfaces that adsorb and retain biomolecular building blocks of the circuit elements (responsible for transferring the charge initiated by electric field or electromagnetic field of light) at the right location, conformation and orientation without loss of functionality. Sophisticated programming of the

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Interfaces in Multiphase Polymer Systems Na+ –COO– –COO– –COO– + Na Na+ –COO– Na+ + – – –COO – Na – –COO –COO –COO + + Na+ Na+ Na Na –COO–

Na+ –COO–

–COO+– Na –COO– + Na – –COO Na+ –COO–

–COOH –COOH –COOH –COOH

–COOH –COOH –COOH –COOH –COOH

105

–COOH –COOH –COOH –COOH

–COOH

–COOH

–

–COO

Insulin

Insulin

pH 4.0

pH 7.4

Glucose oxydase enzyme

Glucose oxydase enzyme

Figure 4.13

Smart insulin releasing pH responsive hydrogel layer.

surfaces is an important aim for designers of tissue engineering scaffolds. The development of smart IPs for hybrid polymer composites and biodegradable polymer systems or even for stabilization (including fire retardancy) is also a challenge for the future.

4.7 Methods of Interface Analysis The surface or interfacial phenomena may relate to various thicknesses (achieving even six orders of magnitude from sub-nanometers to micro- or even millimeters) depending on the interaction considered. For example, wetting behavior is influenced by the composition of the first atomic layers; the optical properties are influenced by interfaces of micrometer thickness, while the adhesion between two polymeric phases modifies the local structure and energy dissipation up to more than one millimeter. Different features of the interfacial region of multiphase systems can be analyzed by methods focusing on different levels and aspects of the structure. The analysis of polymer surfaces and interfaces often require a combination of different techniques (Figure 4.14). By combining the results provided by an appropriately selected set of methods, a complex picture can be formed about the structure and dynamics in the IP. A clear distinction between surface and interface analytical methods can be hardly drawn, therefore we discuss them together in this section. The basic features of the main methods used for surface/interface/interphase analyses are given in Table 4.2 covering different features and thickness ranges of interfacial layers. INTERFACE/INTERPHASE analysis

micromechanical methods

energy, charge analyses

PO MB. SFFT AE

Figure 4.14

CA JKR, W PDT IGC Zeta

mobility detection

microscopic methods

DMA NI TSC

POM TEM SEM SNOM AFM SThM

spectromerty SIMS XPS, AES HREELS SP ELLI RAMAN IR-ATR ssNMR

X ray source

nuclear source

ion beam techniques

PICS~ SANS XCFS~ NR CXD PLS NEXAFS XPEEM TXM

Classification of the main types of surface/interface/interphase analytical methods.

FIB RBS NRA ERD

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Table 4.2 Methods for analyzing the phase border regions. Information specific to interface

Method

Acronym

Technique

Dynamic contact angle

CA

Forces at interfaces Pendant drop

JKR, W PDT

Inverse GC Zeta-potential

IGC ζ potential

Pull-out push-out tests

PO

hydrophobcity resolution ∼ few mm tip/plate moving surface energies drop profile interfacial tension of drop in a liquid substrate adsorbent surface energy of solids electroosmotic electrostatic potential at the boundary fiber/matrix debonding bond strength values, interface and sliding shear strength, fiber/matrix (F/M) interface degradation

Microbond test Single fiber fragmentation Acoustic emission

MB SFFT AE

Dynamic mechanical analysis

DMA

Nanoindentation Thermally stimulated depolarization current

NI TSC

Polarization microscopy

POM

Transmission electron microscopy

TEM

Scanning electron microscopy Scanning near field optical microscopy Atomic force microscopy

SEM SNOM AFM

Scanning thermal microscope Secondary ion mass spectrometry

SThM

X-ray photoelectron spect. Auger spectrometry

XPS, AES

High resolution electron energy loss spectrometry

HREELS

Surface Plasmon spectrometry

SP

Ellipsometry

ELLI

SIMS

wetting by drops

acoustic events

observation of the failure process thermal scanning interaction, damping, vibration discontinuities loss factor, modulus, adhesion local vibration local mechanical properties thermal release of energy states at the interface, frozen polarization loss peaks, Maxwell-Wagner effect hot stage transcrystallization resolution: <10 μm staining structure and density of the interphase lateral and depth resolution 1–3 nm conducting coating surface image, topography, e.g. gold lateral resolution 5 nm confocal, fiber optics, surface topography, lateral AFM resolution: 100 nm tip sensing topography, depth resolution 0.5 nm heated tip sensing local differences of heat conductivity and softness Ga gun in TOF-SIMS IV surface compositions and functions, depth- profiles, resolutions 50 nm energy of electrons, surface composition, depthAuger electrons profiling: 1 nm, lateral resolution: 1 μm inelastic scattering of surface structure, vibration, electrons epitaxial growth, depth resolution ∼1 nm light-surface plasmon 0.1 nm depth and 300 nm interaction lateral resolution on metallic surface Refractive index depth resolution 1 nm

Ref. 156 156, 157 158 159 156, 160 103, 161

162 163, 164 165, 166 167

168 54, 169

170 171

172 173 174 175, 176 177

178

179

180

181

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Table 4.2 (Continued) Method

Acronym

Technique

Micro-Raman

RAMAN

laser spectra

Attenuated total reflectance infrared spect.

IR-ATR, FTIR

diamond sensor surface contact

Solid-state NMR

SSNMR

1

X-ray photon intensity correlation spectroscopy X-ray fluorescence correlation spectroscopy Coherent X-ray diffraction

PICS

coherent beams, fluctuation does not require coherent beams sample smaller than the coherence

Quantitative X-ray microscopy Photoelectron emission microscopy, transmission X-Ray microscopy Small-angle neutron scattering

XFCS CXD

NEXAFS XPEEM TXM SANS

H-13 C MAS, 2D 1 H–13 C WISE

soft X-ray linear dichroism synchrotron source, soft X-ray emission absorption deep penetrating power of neutrons

Neutron reflectivity

NR

diffraction normal to the surface

Positron lifetime spectrometry

PLS

annihilation characteristics

Focused ion beam

FIB

ion accelerator

Rutherford backscattering

RBS

elastic collision

Nuclear Reaction Analysis

NRA

ions of ∼4.4 MeV

Elastic recoil detection, (forward recoil spectr.)

ERD FRES

ion beam analysis, light ion incident

Information specific to interface

Ref.

chemical, physical state mapping 182– 184 kinetic behavior in the vicinity of 185 a surface, depth resolution: 1 μm interphase thickness – spin 186, diffusion, microdomain 187 structure and dynamics evolution of crystal–amorphous 188 interfaces particle dynamics in the 189 interfacial region 3D image of internal strain, 190 coherence: 10 μm transverse, 1 μm longitudinal quantitative, resolution: ∼50 nm, 191 low beam damage, orientation imaging of chemical composition 192 with lateral resolution of ∼30 nm direct determination of chain 193 conformation, depth resolution: 1 nm 1 H, 2 H depth profile, resolution 194 1nm, real-time reflectivity – dynamic proc. free volume distribution, phase 195, transitions, structure secondary 196 bonds imaging, with SEM 197, nanopatterning 198 temperature-dependent 199 interfacial width, interdiffusion IP thickness, contrast by 200 deuteration 201 diffusion, interfaces, depth profiling resolution 5-80 nm contrast: 1 H, 2 H

The determination of the surface and interfacial energies (e.g. surface tension), can be performed by a large variety of methods, such as contact angle measurement, which contributes to the prediction of wetting (which is also influenced by kinetic factors) and even adhesion (if it is dominated by secondary bonds). The result refers to the surface region of typically 0.2 nm and depends on both the surface roughness and contamination. The forces occurring at interfacial contacts can be measured by immersion of

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PDT ZETA JKR CA +++ − − −

Figure 4.15

Scheme of measurement methods relating to energy/charge of a surface.

Wilhelmy-plate (W) method or by detecting the adhesive interaction between a coated scanning probe tip and a well-defined surface according to the Johnson-Kendall-Roberts (JKR) model of elastic adhesive contacts [202]. Solid particles can be used as adsorbent allowing the determination of their surface energy by the inverse gas chromatography (IGC) method. However, in case of nanoparticles this method is hardly applicable. The direct determination of interfacial tension between components of multiphase systems is mostly quite difficult due to high viscosity of components, long time scales for achievement of equilibrium and risk for sample decomposition. The electro-kinetic or zeta potential measurement applies an external electrical field for initiating a relative movement within the electric double layer formed between solid surface and polar liquid phases, which generates an electrical potential or produces an electrical current. The results are influenced by the dissociation of surface groups, the adsorption/substitution of cations, anions or polyelectrolytes and the accumulation or depletion of electrons at the interfaces. The schemes of main relevant tests are assembled in Figure 4.15. The micromechanical tests serve as a bridge between the macroscopic properties and the interfacial characteristics. These are used mainly for analyzing the interfacial contacts in blends and composites. These micromechanical results reflect the influences of interfacial modifications directly, while the macroscopic properties indirectly because of the local variations in voids, concentrations, residual stresses and failures, which cover the interfacial effects partially. The schemes of main micromechanical tests are assembled in Figure 4.16. The mobility of segments and molecules (and the subsequent energy loss of friction), are determined by dynamic mechanical analysis (DMA) and dielectric spectroscopy (such as TSC) and comparison of their results is characteristic to the interfacial interaction. Similar information but localized to the IP region

MB SFFT

PO

Figure 4.16

AE

Scheme of micromechanical tests.

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NI

TSC

DMA

I

+ + +

Figure 4.17

109

+

+

+

+ +

+

+

+ +

+

+ +

+ +

+

+ +

+ + +

+ +

Scheme of methods for determining the interfacial mobility.

is provided when nanoindentation is performed by the AFM system. The schemes of main methods for analyzing the macromolecular mobility are assembled in Figure 4.17. The microscopic methods visualize the IP structure but without identification of the apparent objects the results can be misleading. Therefore the microscopic methods are combined by X-ray analysis (e.g. SEM-EDX) or spectrometry (e.g. micro-Raman). The scanning near-field optical microscopy (SNOM) provides spectroscopic information by illuminating the sample with a tiny light-guiding tip. The atomic force microscopy (AFM, scanning force microscopy) applies several tip-related techniques (hard/soft tapping, phase, friction or chemical contrast) for obtaining information (such as local surface molecular composition) additionally to the topographic image. For chemical recognition and improved resolution the tip can be modified by attachment of specific molecules or nanoparticles (such as carbon nanotubes). Thermal mapping of the sample can be performed by replacing the tip with a thermally-controlled thin platinum wire for the method called Scanning Thermal Microscope (SThM or micro thermal analyser). The analysis of the thermo-mechanical transition at selected positions allows distinguishing phases from each other in multiphase systems [87, 203]. The heated wire-tip can be also used as a nano-manipulation tool. The methods of spectrometry makes feasible the acquisition of micro-structural information of high chemical and/or physical content, monitoring of the physical-chemical interactions and mapping of the local stresses and substances (e.g. IR-ATR, micro-Raman). The SIMS method uses bombardment with an ion beam and analyzes the generated secondary ions in a mass spectrometer. The material can be continuously sputtered away at higher incident ion flux (destructive, dynamic mode) and one can obtain a depth profile of interfaces with a resolution as small as 12 nm depending on the sputter rate. XPS surface analysis determines the energy of electrons removed from the surface atoms by means of X ray beam. It is used to analyze chemical changes at surfaces/interfaces of ∼10 nm induced by thermal or other treatments. High Resolution Electron Energy Loss Spectrometry (HREELS) analyzes a surface layer of less thickness than XPS using ultra high vacuum and incident electrons much smaller than 102 eV. The energy of the scattered electrons (Es ) is measured and the energy loss can be calculated. These electrons have a mean free path of around 1 nm (a few monolayers) in the analyzed materials, which is decreasing with lower energies. The determined intensity versus energy loss diagram allows conclusions to be drawn about surface properties of the sample. Surface Plasmon Spectrometry (SP) detects the thickness and refraction index of a polymer film attached to a thin metal layer on which surface plasmons are excited by the incident light. The interaction between the phases influences the plasmon resonance. Comparable information is provided by ellipsometry (ELLI), which detects the change of the polarization state of reflected (or transmitted) light from a surface. Ellipsometers use oblique incident light within the spectral ranges of interest. The method determines the thickness and refraction index with a resolution better than 1 nm (depending on optical contrast). Liquid cell allows in situ measurement of adsorption kinetics, while information on anisotropy, roughness and variation of thickness or refraction index is provided by ellipsometric imaging. Fourier transform infrared spectroscopy (FTIR) has been used

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Figure 4.18 Signals of methods used for interface analysis (see full names in Table 4.2). (T-temperature, Phphoton, -electron, ⊕-positron, X-x-ray, N-neutron, I-ion, F-force, M-magnetic field, G”-loss modulus, S-sound, C-electric current)

to evaluate the thermal stability and calculate the density of the interphases in polymer composite systems [204, 205]. The Micro-Raman system excites the sample with monochromatic light focused by a microscope objective and analyzes, after filtering the scattered light of the same wavelength, the spectra of the light originating from the inelastic collision with the material. The in-depth and lateral resolution of the method is enhanced when it is combined with SNOM or AFM methods. Solid-state NMR (ssNMR) spectroscopy is a powerful method for characterizing the structure, dynamics and microdomain structure of multiphase polymers [206, 207]. Quantitative determination of the interphase thickness between two immiscible polymers (unsaturated polyester resin and PEO–PPO–PEO triblock copolymer) has been performed for example by 1H spin diffusion experiments [207]. The main methods for surface/interface analysis with their input and output signals are summarized in Figure 4.18. The methods based on synchrotron X ray or neutron, positron sources enhance the resolution compared with their conventional counterparts (such as XPS, IR ellipsometry) and these are applicable not only for determining the local molecular structure at the IP in-static state but also for detecting the dynamic interfacial processes. The X-Ray Fluorescence Correlation Spectroscopy (XFCS) uses a microfocused beam through which particles move in and out, making this method suitable for studying particle dynamics in liquids and solids. This way the interdiffusion of atoms at interfaces between two species of materials; the mobility of the first few molecular layers of a liquid close to a solid surface; the motion of biological macromolecules; or the dynamics of phase separations in liquid, using tracer particles of colloid gold, can be studied. The Positron Lifetime Spectrometry (PLS) provides information about the amorphous phase of polymers by scanning the free volume with positronium atom (the bound state of an electron and a positron). Positron will cease to exist by annihilation when it gets in the vicinity of an electron of the sample forming gamma photons that can

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be detected. The probability of this annihilation process depends on the local density; therefore the fraction of positrons that have a relatively long lifetime gives an insight into the free volume (voids, defects, etc.) structure of the amorphous polymers. Thus specific knowledge can be gained on phase transitions, structural ordering, and the formation of secondary chemical bonds. Ion beam techniques are used either for direct analysis or for assisting other techniques. Focused Ion Beam (FIB) systems use Ga+ ion beams to raster over the surface of a sample in a similar way as the electron beam in a scanning electron microscope. The generated secondary electrons (or ions) are collected to form an image of the surface of the sample. It is mainly used in electron microscope for the creation of a cut perpendicular to the IP allowing the determination of its thickness with good accuracy by cross-sectional imaging. Applications of FIB also includes the preparation for physical-chemical analysis [208]. Rutherford backscattering spectrometry (RBS) means backscattering of a beam of high E using heavy atoms for contrasting the sample. Temperature-dependent interface formation and diffusion between two partially miscible polymers can be analyzed in this way. The optimal temperature for the interaction and the relating maximal achievable interfacial width can be also determined. Nuclear Reaction Analysis (NRA) also uses ions as probes for interface analysis but in this case the deuteration of one component generates contrast between the phases. The target element undergoes a nuclear reaction under resonance conditions for sharply defined resonance energy. To obtain depth information the initial kinetic energy of the ion and its stopping power (energy loss per distance traveled) in the sample is determined (the higher the energy, the deeper the reaction). The methods, using light ions at MeV for elastic collision and analyzing the nucleus recoils and ejects from the target by elastic recoil detection or forward recoil spectroscopy (ERD or FRES), form an own category [209]. The emission angle elastic recoil detection (EA-ERD) benefits from a superb depth resolution, coincidence detection ERD (CD-ERD) has a superior sensitivity (below ppm), while the ˚ and electromagnetic filter detection ERD (E and B ERD) benefits from a depth resolution of ca. 100 A sensitivity better than ≈0.2 at.%.

4.8 Conclusions Interfaces determine the macroscopic behavior of multiphase materials consisting of polymer, metal, ceramic, glass, natural and biomaterial components and influence their applicability in traditional and advanced areas. Their research represents an essential part of material science. The advancement of a fundamental understanding of the structures of phase borders and their role in controlling or influencing fracture, creep, superplasticity, charge transfer and storage, magnetic properties, corrosion, flammability, biological response, etc., accelerates the development of the whole science and technology. This importance became even more evident since the nanostructured materials have begun to dominate most of the key research areas, such as semiconductors, drug delivery, biomaterials and fire retardancy. The tremendous advancement in theory, simulation, specified experimental techniques and instrumentation, generated recently in the science of interfaces, initiated the intention of this chapter to give a comprehensive discussion and classification of the relevant issues of this complex area. The provided systematic classification, considering static and responsive structural, physical, chemical biological and analytical aspects of interphases, tries to contribute to the renewal of the relevant technologies. The wide scope did not exclude the consideration of the common thermodynamic and kinetic fundamentals and the evaluation how these basic features establish the macroscopic properties of the multiphase systems. By promoting the understanding of these main characteristics a basis is provided for elaborating strategy for precise design and engineering of optimal interfacial structures. The discussion of the different ways of interfacial modification included new aspects among the known topics. Beyond the transcrystalline interlayer the importance of amorphous interphase in the fields of

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ceramics and pharmaceuticals is highlighted. The combined discussion of the topics being in relation to the modification of segmental mobility at the interfaces (a diffusion, permittivity and compliance) gives a novel view for their further development. Among the physical and reactive ways for improving the interfacial connection between hardly compatible phases special methods are discussed that were missing from the former reviews, such as nanocompatibilization, reactive surfactants and hierarchic multilayer interfacial structures (mimicking nature). New perspectives of living polymerization and in-line electron beam treatment are introduced beyond the traditional methods of reactive compatibilization. The catalytic activity of interfaces can be disadvantageous in the case of degradation but advantageous when interfacial polymerization or controlled fire retardancy is required. The insulating and activating interlayers, developed for regulating these reactions, are also discussed. Newly-developed responsive interphases, appearing in the materials developed recently, enable to combine the sensing with the common functions of static ones. The use of non-reversible adaptive variants is shown in the field of fire retardancy, while the rapid spread of the smart types is demonstrated through the biomedical application. Development of the interfacial modification requires parallel advancement in the analysis of interphases. The large variation of the aims involves wide scale of investigations, the common discussion of which means a big challenge. The classification provided here involves the main techniques some of which were omitted by the former reviews. Simple figures were constructed for promoting better understanding. Believing that the rapid development of the engineered interphases in multiphase systems is just at the beginning it is hoped that this wide scale discussion initiate emergence of new ideas for industrial application even at areas not covered here.

References 1. K.E. Geckeler, F. Rupp., J. Geis-Gerstorfer, Interfaces and interphases of (Bio)materials: Definitions, structures, and dynamics. Adv. Mater. 9, 513–518 (1997). 2. J. Karger-Kocsis, O. Gryshchuk, S. Schmitt, Vinylester/epoxy-based thermosets of interpenetrating network structure: an atomic force microscopic study. J Mater Sci., 38, 413–20. (2003). 3. J. Karger-Kocsis, O. Gryshchuk, N. Jost, Toughness response of vinylester/epoxy-based thermosets of interpenetrating network structure as a function of the epoxy resin formulation: effects of the cyclohexylene linkage. J. Appl. Polym. Sci., 88, 2124–31. (2003). 4. I. Chodak, I. Novak, (in Polypropylene, An A-Z Reference), J. Karger-Kocsis (Ed), Kluwer Academic Publisher, Dordrecht/Boston/London, 1999, p. 790. 5. M. Geoghegan, G. Krausch, Wetting at polymer surfaces and interfaces, Progress in Polymer Science 28(2): 261–302 (2003). 6. D. Cunliffe, S. Pennadam, C. Alexander, Synthetic and biological polymers-merging the interface, European Polymer Journal, 40(1), 5–25 (2004). 7. R. J. Kroeze M. N. Helder, L. E. Govaert, T. H. Smit, Biodegradable polymers in bone tissue, Engineering Materials, 2, 833–856; (2009). 8. W.V. Gutowski. Interface/interphase engineering of polymers for adhesion enhancement: Review of micromechanical aspects of polymer interface reinforcement through surface grafted molecular brushes, J. Adhesion; 79, 445–482 (2003). 9. U. Landman, W.D. Luedtke, Nanomechanics and dynamics of tip substrate interactions, J. Vacuum Sci. Technol.; B9 414–423 (1991). 10. C. Zhu, G. Bao, N. Wang, Cell mechanics: Mechanical response, cell adhesion, and molecular deformation. Ann. Rev. Biomed. Eng. 2, 189–226 (2000). 11. K. R. Shull, Contact mechanics and the adhesion of soft solids, Mater. Sci. Eng. R. 36, 1–45 (2002). 12. S. N. Jamadagni, R. Godawat, J. S. Dordick, S. Garde, How interfaces affect hydrophobically driven polymer folding, Journal of Physical Chemistry B,. 113(13) 4093–4101 (2009).

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13. Y.S. Lipatov, Interfaces in polymer–polymer composites. In: Controlled interfaces in composite materials, Ishida H. (Ed.), Elsevier, Amsterdam; 1990 pp. 599–612. 14. R. A. L. Jones, R. W. Richards, In: Polymers at surfaces and interfaces. Cambridge University Press, Cambridge, 1999. 15. W. V. Gutowski, S. Li, L. Russell, C. Filippou, P. Hoobin, S. Petinakis, Effect of surface-grafted molecular brushes on the adhesion performance of bonded polymers and composite interfaces, Composite Interfaces, 9(1): 89–133 (2002). 16. Y.S. Lipatov, Polymer blends and interpenetrating polymer networks at the interface with solids, Progress in Polymer Science, 27(9) 1721–1801 (2002). 17. A. E. Nesterov, Y. S. Lipatov, Compatibilizing effect of a filler in binary polymer mixtures, Polymer; 40, 1347–9. (1998). 18. M. Si, V. Zaitsev, M. Goldman, A. Frenkel, D. G. Peiffer, E. Weil, J. C. Sokolov, M. H. Rafailovich, Self-extinguishing polymer/organoclay nanocomposites Polymer Deg. Stab., 92, 86–93 (2007). 19. A. R. Khokhlov, Influence of excluded volume effect on the rates of chemically controlled polymer-polymer reactions Makromolekulare chemie-rapid communications 2, 633 (1981). 20. B. K. Finka, R. L. McCullough, Interphase research issues, Composites: Part A; 30, 1–2 (1999). 21. W. S. Gutowski, Physical model of interface and interphase performance in composite materials and bonded polymers, Angew. Makromol. Chem., 272, 51–56 (1999). 22. O. Okay, W. Oppermann, Polyacrylamide−Clay Nanocomposite Hydrogels: Rheological and Light Scattering Characterization, Macromolecules, 40(9), 3378–3387 (2007). 23. B. C. Ray, Temperature Effect During Humid Ageing on Interfaces of Glass and Carbon Fibers Reinforced Epoxy Composites, Journal of Colloid and Interface Science 298, 111–117 (2006). 24. G. Alawsib, S. Aldajaha, S. A. Rahmaan, Impact of humidity on the durability of E-glass/polymer composites, Materials & Design 30(7), 2506–2512 (2009) 25. F. A. Ramirez., L. A. Carlsson, B. A Acha, Evaluation of water degradation of vinylester and epoxy matrix composites by single fiber and composite tests. J Mater Sci. 43(15) 5230–5242 (2008) 26. A. Kelly, W. R. Tyson, Tensile properties of fibre-reinforced metals, J. Mech. Phys. Solids 13, 329 (1965). 27. B. Puk´anszky, Interfaces and interphases in multicomponent materials: past, present, future, European Polymer Journal 41(4), 645–662 (2005). 28. Gy. Voros, B. Pukanszky, Modeling of the effect of a soft interlayer on the stress distribution around fibers: longitudinal and transverse loading. Macromol. Mater. Eng,; 287, 139–48 (2002). 29. N. Matsumoto and J. A. Nairn, The fracture toughness of medium density fiberboard (MDF) including the effects of fiber bridging and crack–plane interference, Engineering Fracture Mechanics, 76, 2748–2757 (2010). 30. J-K. Kim, Y-W. Mai, In: Engineered interfaces in fiber reinforced composites, Ch 3. Elsevier, Amsterdam, 1998, p. 43. 31. A.T. DiBenedetto, Tailoring of interfaces in glass fiber reinforced polymer composites: A review, Mater. Sci. Eng., A 302, 74–82 (2001). 32. J. Jancar, Review of the role of the interphase in the control of composite performance on micro- and nano-length scales, J. Materials Sci., 43, 6747–6757 (2008). 33. Y. P. Chen, Atomistic viewpoint of the applicability of microcontinuum theories, International J. Solids and Structures, 41, 2085 (2004). 34. W. Wang and A. T. DiBenedetto, A modified silane treatment for superior stabilityof glass reinforced composites, J. Adhesion, 68, 183 (1999). 35. C. Galiotis, in Multi-scale Modelling of Composite Material Systems, P. W. R. Beaumont, C. Soutis (Eds) Woodhead Publ. Inc., Cambridge (2005), p. 33. 36. P. Sharma, S, Ganti N. Bhate, Effect of surfaces on the size-dependent elastic state of nano-inhomogeneities, Appl Phys Lett., 82, 535–537 (2003). 37. Y. Chen, J. D. Lee, A. Eskandarian, Examining the physical foundation of continuum theories from viewpoint of phonon dispersion relations, Intern. J. Eng. Sci., 41, 61–83 (2003). 38. B. K. Fink, R. L. McCullough, Interphase research issues, Composites Part A: Applied Science and Manufacturing, 30(1), 1–2 (1999).

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39. K. Okada, K. Watanabe, T. Urushihara, A. Toda and M. Hikoshaka, Role of epitaxy of nucleating agent (NA) in nucleation mechanism of polymers, Polymer, 48, 401–408 (2007). 40. Y. C. Chua, X. Lu, Polymorphism, behavior of poly(ethylene naphthalate)/clay nanocomposites: Role of clay surface modification, Langmuir, 23, 1701–1710, (2007). 41. J. C. Wittmann, B. Lotz, Epitaxial crystallization of polymers on organic and polymeric substrates, Progr. in Polym. Sci., 15(6), 909–948. (1990). 42. A. R. Sanadi, D. F. Caulfield, Transcrystalline interphases in natural fiber–PP composites: Effect of coupling agent, Composite Interfaces, 7, 31–43 (2000). 43. S.V. Bobylev, I. A. Ovid’ko, Dislocation nucleation at amorphous intergrain boundaries in deformed nanoceramics, Physics of the Solid State, 50(4), 642–648 (2008). 44. B. Vajna, Gy. Marosi, I. Farkas, Zs. Zsigmond, Raman microscopic characterization of the structure of tablets manufactured by different methods, J. Pharm. Biomed. Anal. 51(1), 30–38 (2010). 45. D. T. Friesen, W. C. Babcock, R. M. Shanker, D. T. Smithey, R. Tadday, Pharmaceutical Compositions of Adsorbates of Amorphous Drug US Patent WO/2003/000238 (2003) 46. S. Rogovina, G. Vikhoreva, Polysaccharide-based polymer blends: Methods of their production, Glycoconjugate J., 23, 611–618 (2006). 47. G. Marosi, P. Anna, S. Nagy, S. Matko, I. Antal, New concepts in formulation form extrusion to quantitative micro-spectroscopy, Eur. J. Pharm. Sci., 25(1), S27–S29 (2005). 48. M.D. Sanchez-Garcia, E. Gimenez, and J. M. Lagaron, Novel PET nanocomposites of interest in food packaging applications and comparative barrier performance with biopolyester nanocomposites, J. Plast. Film Sheet., 23, 133–148 (2007). 49. Y. Pu, M. H. Rafailovich, J. Sokolov, D. Gersappe, T. Peterson, W. L. Wu, S. A. Schwarz, Mobility of polymer chains confined at a free surface, Phys. Rev. Lett., 87, 6101 (2001). 50. X. Zheng, M. H. Rafailovich, J. Sokolov, Y. Strzhemechny, S. A. Schwarz, B. B. Sauer, M. Rubinstein, Long-range effects on polymer diffusion induced by a bounding interface, Phys. Rev. Lett. 79, 241–244 (1997). 51. S. Granick, S. K. Kumar, E. J. Amis, M. Antonietti, A. C. Balazs, A. K. Chakraborty, G. S. Grest, C. Hawker, P. Janmey, E. J. Kramer, R. Nuzzo, T. P. Russell, C. R. Safinya, Macromolecules at surfaces: Research challenges and opportunities from tribology to bBiology, J. Polym. Sci. Part B: Polym. Phys., 41, 2755–2818 (2003). 52. X. J. Zhang,. S. A Jenekhe, Electroluminescence of multicomponent conjugated polymers. 1. Roles of polymer/polymer interfaces in emission enhancement and voltage-tunable multicolor emission in semiconducting polymer/polymer heterojunctions, Macromolecules, 33, 2069–2082 (2000). 53. P. Tewari, E. Furman, R. Rajagopalan, M. T. Lanagan, T. Michael, Electrical insulation, IEEE International Symposium on Electrical Insulation, Conference Record, 312–315 (2008). 54. M. Kuliev, R. Ismaiilova, M. Bairamov, Study of polymer composites by methods of thermoactivated spectroscopy, Surf. Eng. Appl. Electrochem., 44(6), 467–470 (2008). 55. H. Lu, X. Zhang, H. Zhang, Influence of the relaxation of Maxwell-Wagner-Sillars polarization and dc conductivity on the dielectric behaviors of nylon 1010, J. Appl. Phys. 100, 054104 (2006). 56. G. Banhegyi, G. Marosi, G. Bertalan, F. E. Karasz, Studies of thermally stimulated current in polypropylene/calcium carbonate/surfactant systems, Colloid Polym. Sci., 270(2), 113–127 (1992). 57. N. G. Green, A. Ramos, H. Morgan, Ac electrokinetics: A survey of sub-micrometre particle dynamics, J. Phys. D: Appl. Phys., 33, 632–641 (2000). 58. B. N. Pal, S. Basu, D Chakravorty, Electrically aligned binary system of nanoparticles, J. Appl. Phys., 98, 84306 (2005). 59. M. J. Casavant, D. A. Walters, J. J. Schmidt, R. E. Smalley, Neat macroscopic membranes of aligned carbon nanotubes, J. Appl. Phys., 93, 2153–56 (2003). 60. C. Park, J. Wilkinson, S. Banda, Z. Ounaies, K. Wise, G. Sauti, P. T. Lillehei, J. Harrison, Aligned single-wall carbon nanotube, J. Polym. Sci., Part B: Polym. Phys. 44, 1751–1762 (2006). 61. H. Koerner, J. D. Jacobs, D. W. Tomlin, J. Busbee, R. Vaia, Polymer-encapsulated liquid crystals comprising montmorillonite clay, Adv. Mater., 16, 297–302 (2004). 62. Y. Cao, P. C. Irwin, K. Younsi, The future of nanodielectrics in the electric power industry, IEEE Trans. Dielectr. Electr. Insul., 11 (5), 797–807 (2004).

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63. M. Roy, J. K. Nelson, R. K. MacCrone, L. S. Schadler, C. W. Rees, R. Keefe, W. Zenger, Polymer nanocomposite dielectrics – the role of the interface, IEEE Trans. Dielectr. Electr. Insul., 12 (4), 629–43 (2005). 64. J. Wang, N. Huang, Studies of the action of silicone at the interphase on the flame retardant system of EBA copolymer-chalk, J. Fire Sci., 27(3), 235–264 (2009). 65. S. Kim, L. T. Drzal, Improvement of electric conductivity of LLDPE based nanocomposite by paraffin coating on exfoliated graphite nanoplatelets, Composites Part A: Applied Science and Manufacturing 2010, in press. 66. H. D. Wagner, R. A. Vaia, Nanocomposites: Issues at the interface. Mater Today, 7(11): 38–42 (2004). 67. J. Moczo, E. Fekete, B. Pukanszky, Acid-base interactions and interphase formation in particulate-filled polymers. J. Adhesion, 78, 861–875 (2002). 68. G. Marosi, A. Tohl, G. Bertalan, P. Anna, M. A. Maatoug, I. Ravadits, I. Bertoti, A. Toth, Modified interfaces in multicomponent polypropylene fibers, Compos. Pt. A-Appl. Sci. Manuf., 29(9–10), 1305–1311 (1998). ¨ 69. A. Buz´agh, Uber den Zusammenhang zwischen Adh¨asion und Stabilit¨at disperser Systeme, Kolloid Z. 125, 14–21 (1952). 70. J. L. Kardos., in Molecular Characterisation of Composites Interfaces, H. Ishida, G. Kumar (Eds), Plenum Publ. Corp. (1985). 71. A. K. Chakraborty, H. T. Davis, M. Tirrell, A molecular orbital study of the interactions of acrylic polymers with aluminum: Implications for adhesion, J. Polym. Sci. A: Polym. Chem., 28, 3185–3195 (1990). 72. G. Marosi, G. Bertalan, I. Rusznak, P. Anna, Role of interfacial layers in the properties of particle-filled polyolefin systems, Colloids and Surfaces, 23(3), 185–198 (1986). 73. M. U. Wahit, A. Hassan, Z. A. Mohd-Ishak, A. R. Rahmat, A. Abu Bakar, Morphology, thermal, and mechanical behavior of ethylene octene copolymer toughened polyamide 6/polypropylene nanocomposites, J. Thermoplast. Comp. Mater., 19(5), 545–567 (2006). 74. Gy. Marosi, R. L´agner, Gy. Bertalan, P. Anna, A. Tohl., Thermoanalytical study of nucleating effects in polypropylene composites II. Filler and elastomer containing polypropylene. J. Therm. Anal., 47, 1163 (1996). 75. B. Ji, H. Gao, Mechanical properties of nanostructure of biological materials J. Mech. Phys. Solids, 52, 1963–1990 (2004). 76. B. Ji, H. Gao, Elastic properties of nanocomposite structure of bone, Comp. Sci. Technol., 66, 1212–1218 (2006). 77. G. Fantner, E. Oroudjev, G. Schitter, L. S. Golde, P. Thurner, M. M. Finch, P. Turner, T. Gutsmann, D. E. Morse, H. Hansma, P. K. Hansma, Sacrificial bonds and hidden length: Unraveling molecular mesostructures in tough materials, Biophys J., 90, 1411–1418 (2006). 78. G. E. Fantner, O. Rabinovych, G. Schitter, P. Thurner, J. H. Kindt, M. M. Finch, J. C. Weaver, L. S. Golde, D. E. Morse, A. Lipman, E. I. W., Rangelow, P. K. Hansma, Compos Sci Technol., 66, 1205 (2006). 79. J. Jagur-Grodzinski, Preparation of functionalized polymers using living and controlled polymerizations, Reactive Functional Polym., 49, 1–54 (2001). 80. J. Ladd, Z. Zhang, S. Chen, J. C. Hower, S. Jiang, Zwitterionic polymers Exhibiting high resistance to nonspecific protein adsorption from human serum and plasma, Biomacromolecules, 9(5), 1357–1361 (2008). 81. R.E. Gorga, K. K. S. Lau, K. K. Gleason, R. E. Cohen. The importance of interfacial design at the carbon nanotube/polymer composite interface, J. Appl. Pol.. Sci. 102(2), 1413–1418 (2006). 82. A. Meyer-Plath, K. Schroder, B. Finke, A. Ohl, Current trends in biomaterial surface functionalization – Nitrogencontaining plasma assisted processes with enhanced selectivity. Vacuum, 71, 391–406 (2003). 83. Y. Yang, S. Basu, D. L. Tomasko, L. J. Lee, S.-T. Yang, Fabrication of well defined PLGA scaffolds using novel micro-embossing and carbon dioxide bonding, Biomaterials, 26(15), 2585–2594 (2005). 84. F. Ciardelli, S. Coiai, E. Passaglia, A. Pucci, G. Ruggeri, Nanocomposites based on polyolefins and functional thermoplastic materials, Polymer International, 57(6), 805–836 (2008). 85. W. S. Gutowski, A. Bilyk, S. Li, M. Espiritu, I. Burgar, The influence of structure of the interface and interphase on paint adhesion, Composite Interfaces, 12(8–9), 817–835 (2005). 86. S. D. Faulkner, Y. W. Kwon, S. Bartlett, E. A. Rasmussen, Study of composite joint strength with carbon nanotube reinforcement, J. Mater. Sci. 44(11), 2858–2864 (2009). 87. A. Szabo, B. Marosfoi, P. Anna, G. Marosi, Complex micro-analysis assisted design of fire retardant nanocomposites – Contribution to the nano-mechanism, in Fire Retardancy of Polymers, New Strategies and Mechanisms, K. Kandola and R. Hull (Eds), The Royal Society of Chemistry 4 (2008), pp. 74–91.

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88. S. Matko, I. Repasi, A. Szabo, B. Bodzay, P. Anna, G. Marosi, Fire retardancy and environmental assessment of rubbery blends of recycled polymers, Express Polym. Lett. 2(2), 126–132 (2008). 89. G. Marosi, G. Bertalan, Role of interfaces in multicomponent polymer systems, in Modification and Blending of Synthetic and Natural Macromolecules for Preparing Multiphase Structure and Functional Materials, NATO Advanced Study Institute Series, F. Ciardelli, S. Penczek (Eds), Pisa, 175 (2004), pp. 135– 154. 90. C. Harrats, G. Groeninckx, Reactive processing of polymer blend using reactive compatibilization and dynamic crosslinking: phase morphology control and microstructure – Property relations, in Modification and Blending of Synthetic and Natural Macromolecules for Preparing Multiphase Structure and Functional Materials, NATO Advanced Study Institute Series, F. Ciardelli, S. Penczek (Eds), Pisa, 175 (2004), pp. 155–200. 91. C. W. Macosko, H. K. Jeon, T. R. Hoye, Reactions at polymer-polymer interfaces for blend compatibilization, Progr. Polym. Sci., 30(8–9), 939–947 (2005). 92. A. T. DiBenedetto, Tailoring of interfaces in glass fiber reinforced polymer composites, Mat. Sci. Eng. A - Struct., 302(1), 74–82 (2001). 93. T. D. Jones, J. S. Schulze, C. W. Macosko, T. P. Lodge, Effect of thermodynamic interactions on reactions at polymer/polymer interfaces, Macromolecules, 36(19), 7212–7219 (2003). 94. V. B. Sadhu, J. Pionteck, P. Potschke, L. Jakisch, A. Janke, Creation of crosslinkable interphases in polymer blends, Macromol. Symp., 210: 165–174 (2004). 95. I. Ravadits, A. oth, G. Marosi, A. Marton, A. Szep, Organosilicon surface layer on polyolefins to achieve improved flame retardancy through oxygen barrier effect. Polym. Degrad. Stab., 74(3), 414–422 (2001). 96. A. Qu´ed´e, C. Jama, P. Supiot, M. Le Bras, R. Delobel, O. Dessaux P. Goudmand, Elaboration of fire retardant coatings on polyamide-6 using a cold plasma polymerization process, Surface and Coatings Technology, 151–152, 424–428 (2002). 97. M. Gehde, U. Gohs, G. Heinrich, K. Michael, K. Sritragool, U. Wagenknecht, S. Wießner, Modifizierung von Gummipartikel/Thermoplastmischungen mit energiereichen Elektronen, Gummi Fasern Kunststoffe 62, 88–93 (2009). 98. Gy. Marosi, P. Anna, I. Csontos, A. M´arton, Gy. Bertalan, New reactive additives for interface modification in multicomponent polyolefin systems, Macromol. Symp., 176, 189–198 (2001). 99. Gy. Bertalan, Gy. Marosi, P. Anna, I. Ravadits, I. Csontos, A. T´oth, Role of interface modification in filled and flame retarded polymer systems, Solid State Ionics, 141–142, 211–215 (2001). 100. Gy. Marosi, P. Anna, A. Marton, I. Csontos, Sz. Matko, A. Szep, E. Kiss Reactive surfactants – New type of additives for polymers, Progr. Coll. Polym. Sci., 125, 189–193 (2004). 101. P. Anna, Gy. Bertalan, Gy. Marosi, I. Ravadits, M.A. Maatoug, Effect of interface modification on the photo-stability of pigmented polyethylene films, Polym. Degrad. Stab. 73, 463–466 (2001). 102. X. F. Zhoua, H. D. Wagner, S. R. Nutt, Interfacial properties of polymer composites measured by push-out and fragmentation tests, Composites Part A: Applied Science and Manufacturing 40, 3, 235–320 (2009). 103. J. Jancar, Engineered interphases in polypropylene composites, in Handbook of Polypropylene and Polypropylene Composites, G. Harutun, G. Karian (Eds), Marcel Dekker New York (2003), pp. 413–464. 104. J. Draper, I. Luzinov, S. Minko, I. Tokarev, M. Stamm, Mixed polymer brushes by sequential polymer addition: Anchoring layer effect, Langmuir, 20(10), 4064–4075 (2004). 105. J. J. Lesko, K. Jayaraman, K. L. Reifsnider, Gradient interphase regions in composite systems, Key Eng. Mat., 116, 61–86 (1996). 106. K. Matyjaszewski, Preparation of nanocomposites with well-defined organic (co)polymers, in Modification and Blending of Synthetic and Natural Macromolecules for Preparing Multiphase Structure and Functional Materials, NATO Advanced Study Institute Series, F. Ciardelli, S. Penczek (Eds) Pisa, 175, (2004), pp. 135–154. 107. S. Edmondson, V. L. Osborne, W. T. S. Huck, Polymer brushes via surface-initiated polymerizations, Chem. Soc. Rev., 33, 14–22 (2004). 108. J. Pyun, S. Jia, T. Kowalewski, K. Matyjaszewski, Synthesis and surface attachment of ABC triblock copolymers containing glassy and rubbery segments, Macromol. Chem. Phys., 205, 411–417 (2004). 109. T. Wu, K. Efimenko, J. Genzer, Preparing high-density polymer brushes by mechanically assisted polymer assembly, Macromolecules, 34(4), 684–686, (2001).

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110. I. Luzinov, S. Minko, V. V. Tsukruk, Adaptive and responsive surfaces through controlled reorganization of interfacial polymer layers, Progr. Polym. Sci. 29(7) (2004), 635–698. 111. M. Krouit, J. Bras, M. N. Belgacem, Cellulose surface grafting with polycaprolactone by heterogeneous clickchemistry, Eur. Polym. J. 44(12), 4074–4081 (2008). 112. R. Partch, S. Brown, Aerosol and solution modification of particle-polymer interfaces, J. of Adhesion, 67(1–4), 259–276 (1998). 113. I. Ruszn´ak, Gy. Bertalan, P. Anna, Gy. Marosi, I. Moln´ar, Wechselwirkung zwischen Polymer und F¨ullstoff in Polymerverbunden, Plaste Kautsch., 32, 254–262 (1985). 114. G. Marosi, G. Bertalan, I. Rusznak, P. Anna, I. Molnar, The effect of additives on the structure and mechanical properties of PP-filler system, Polymer Composites, B. Sedlacek (Ed.), Walter de Gruyter, Berlin (1986), pp. 457–464. 115. Gy. Marosi, Gy. Bertalan Gy, P. Anna, I. Ruszn´ak, Elastomer interphase in particle filled polypropylene, structure, formation and mechanical characteristics, J. Polymer Eng., 12:34, (1993). 116. J. Jancar, Interphase phenomena in Nano- and Micro-Mechanics of Ppolymer Bblends and Composites, J. KargerKocsis, S. Fakirov (Eds), Hanser (2009), p. 241. 117. P. Crowley, L. G. Martini, Excipients for pharmaceutical dosage forms, in Encyclopedia of Pharmaceutical Technology, J. Swarbrick (Ed.), 3 Informa Healthcare New York (2007), p. 1609. 118. M. A. Maatoug, P. Anna, G. Bertalan, I. Ravadits, G. Marosi, I. Csontos, A. Marton, A. Toth, Role of pigments in the stability of polyethylene systems, Macromol. Mater. Eng., 282(9), 30–36 (2000). 119. Gy. Marosi, P. Anna, A. Marton, Gy. Bertalan, A. Bota, A. Toth, M. Mohai, I. Racz, Flame retarded polyolefin systems of controlled interphase, Polym. Adv. Technol., 13(10–12), 1103–1111 (2002). 120. M. Jimenez, S. Duquesne, S. Bourbigot, Intumescent fire protective coating: Toward a better understanding of their mechanism of action, Thermochimica Acta, 449(1–2), 16–26 (2006). 121. P. Anna, G. Marosi, I. Csontos, S. Bourbigot, M. Le Bras, R. Delobel, Influence of modified rheology on the efficiency of intumescent flame retardant systems, Polymer Degradation and Stability, 74, 423–426 (2001). 122. F. Peter, C. Zarcula, C. Kiss, C. Csunderlik, L. Poppe, Enhancement of lipases enantioselectivity by entrapment in hydrophobic sol-gel materials: Influence of silane precursors and immobilization parameters, J. Biotechnology, 131(2), S109 (2007). 123. G. Marosi, G. Bertalan, P. Anna, I. Ravadits, S. Bourbigot, M. Le Bras, R. Delobel, Polyphosphate flame retardants in polyolefins, in Recent Advances in Flame Retardancy of Polymeric Materials XI, Lewin, M. (Ed.) N BCC. Inc. (2000), pp. 154–63. 124. S. Bourbigot, M. Le Bras, S. Duquesne, M. Rochery, Recent advances for intumescent polymers, Macromolecular Materials and Engineering, 289, 499–511 (2004). 125. S. Keszei, Sz. Matk´o, Gy. Bertalan, P. Anna, Gy. Marosi, A. T´oth, Progress in interface modifications: from compatibilization to adaptive and smart interphases, Eur. Polym. J., 41, 697–705 (2005). 126. M. L. Hans, A. M. Lowman, Biodegradable nanoparticles for drug delivery and targeting, Curr. Opin. Solid. St. Mat., 6(4): 319–327 (2002). 127. K. A. Mowery, M. H. Schoenfisch, J. E. Saavedra, L. K. Keefer, M. E. Meyerhoff, Preparation and characterization of hydrophobic polymeric films that are thromboresistant via nitric oxide release, Biomaterials, 21, 9–21 (2000). 128. N. Billingham, Does additive migration still matter? J. Vinyl. Addit. Techn. 9(3), 105–105 (2003). 129. Gy. Marosi, B. B. Marosf¨oi, B. Bodzay, T. Igricz, Gy. Bertalan, Design of interlayers for fire retarded polymeric systems, in Fire Retardancy of Polymeric Materials, C. A. Wilkie, A. B. Morgan (Eds), Taylor and Francis CRC Press, London (2010), pp. 329–348. 130. Gy. Marosi, P. Anna, Gy. Bertalan, Sz. Szab´o, I. Ravadits, J. Papp, Role of interface modification in flame retarded multiphase polyolefin systems, in Fire and Polymers, Nelson, G., Wilkie, C. (Eds), ACS Ser., 797 Washington (2001), pp. 161–171. 131. Gy. Marosi, I. Ravadits, Gy. Bertalan, P. Anna, M.A. Maatoug, A. T´oth, M. D. Tran, Role of migration process in the efficiency of intumescent flame retardant additives, in Polypropylene in Fire Retardancy of Polymers: the Use of Intumescence, Camino G., Le Bras, M., Bourbigot, S., Delobel, R. (Eds), Royal Soc. of Chem. Cambridge (1998), p. 325.

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132. G. Marosi, A. Marton, A. Szep, I. Csontos, S. Keszei, E. Zimonyi, A. Toth, X. Almeras, Fire retardancy effect of migration in polypropylene nanocomposites induced by modified interlayer, Polym. Degrad. Stabil., 82(2), 379–385 (2003). 133. G. Marosi, A. Marton, P. Anna, G. Bertalan, B. Marosfoi, A. Szep, Ceramic precursor in flame retardant systems, Polym. Degrad. Stab., 77(2), 259–265 (2002). 134. B. Marosfoi, A. Szabo, K. Kiss, G. Marosi, Use of organosilicone composites as flame retardant additives and coating for polypropylene, in Fire Retardancy of Polymers, New Strategies and Mechanisms, K. Kandola and R. Hull (Eds), The Royal Society of Chemistry 4 (2008), pp. 49–58. 135. G. Speranza, G. Gottardi, C. Pederzolli, Role of chemical interactions in bacterial adhesion to polymer surfaces, Biomaterials; 25(11), 2029–2037 (2004). 136. B. D. Ratner, Molecular design strategies for biomaterials that heal, Macromol. Sci., 30, 327–335 (1998). 137. M. Rousseau, E. Lopez, P. Stempfl´e, M. Brendl´e, L. Franke, A. Guette, R. Naslain, X. Bourrat, Multiscale structure of sheet nacre, Biomaterials, 26(31), 6254–6262 (2005). 138. M. L. Hans, A. M. Lowman, Biodegradable nanoparticles for drug delivery and targeting, Curr. Opin. Solid. St. M., 6(4), 319–327 (2002). 139. K. A. Mowery, M. H. Schoenfisch, J. E. Saavedra, L. K. Keefer, M. E. Meyerhoff, Preparation and characterization of hydrophobic polymeric films that are thromboresistant via nitric oxide release, Biomaterials, 21, 9–21 (2000). 140. J. S. Chena, M. D. Poliksa, C. K. Ober, Y. Zhanga, U. Wiesnera, E. Giannelis, Study of the interlayer expansion mechanism and thermal–mechanical properties of surface-initiated epoxy nanocomposites, Polymer 43(18), 4895–4904 (2002). 141. A. Toldy, P. Anna, I. Csontos, A. Szabo, G. Marosi, Intrinsically flame retardant epoxy resin – Fire performance and background - Part I, Polym. Degrad. Stab., 92, 2223–30 (2007). 142. Z. Guo, Z. Fang, L. Tong, Application of percolation model on the brittle to ductile transition for polystyrene and polyolefin elastomer blends. Exp. Polym. Lett. 1(1) 37–43 (2007). 143. F. Zhang, J. Zhang, Y. Wang, Modeling study on the combustion of intumescent fire-retardant polypropylene. Exp. Polym. Lett., 1(3), 157–65 (2007). 144. G. Marosi, P. Anna, A. M´arton, S. Matk´o, A. Szep, S. Keszei, I. Csontos, B. Marosf¨oi, Mechanism of interactions in flame retarded polymer nanocomposites, in International Conference on Additives, Wilkie, C. and Al-Malaika, S (Eds), San Francisco: ECM Ltd. 12, 203, (2003) 145. M. Tirrell, E. Kokkoli, M. Biesalski, The role of surface science in bioengineered materials Surface Sci., 500, 61–83 (2002). 146. M. Zr´ınyi, D. Szab´o, Muscular contraction mimicked by magnetic gels, Int. J. Mod. Phys., B15(6–7), 557–563 (2001). 147. D. G. Castner, B. D. Ratner, Biomedical surface science: Foundations to frontiers. Surf. Sci., 500(1–3), 28–60 (2002). 148. N. Nath, A. Chilkoti, Creating “smart” surfaces using stimuli responsive polymers. Adv. Mater., 14(17), 1243 (2002). 149. J.-H. Kim, T. R. Lee, Thermo- and pH-responsive hydrogel-coated gold nanoparticles, Chem. Mater., 16(19), 3647–3651 (2004). 150. X. Zeng, H. Jiang, Tunable liquid microlens actuated by infrared light-responsive hydrogel, Appl. Phys. Lett., 93, 151101 (2008). 151. I. Roy, M. N. Gupta, Smart polymeric materials emerging biochemical applications, Chem. Biol., 10(12): 1161–1171 (2003). 152. A. Kikuchi, T. Okano, Intelligent thermoresponsive polymeric stationary phases for aqueous chromatography of biological compounds, Progr. Polym Sci., 27(6), 1165–1193 (2002). 153. N. A. Peppas, C. D. Bures, Glucose-responsive hydrogels, in Encyclopedia of Biomater. Biomed. Eng., 2, G. E. Wnek, G. L. Bowlin (Eds), Informa-Healthcare, New York (2008), p. 1163. 154. N. Tsubokawa, Functionalization of carbon material by surface grafting of polymers, Bull. Chem. Soc. Jap. 75(10), 2115–2136 (2002). 155. S. H. Anastasiadis, H. Retsos, S. Pispas, N. Hadjichristidis, S. Neophytides, Smart polymer surfaces, Macromolecules; 36(6), 1994–1999 (2003).

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156. K. Grundke, Characterization of polymer surfaces by wetting and electrokinetic measurements – Contact angle, interfacial tension, zeta potential, in Polymer Surfaces and Interfaces Characterization, Modification and Applications, M. Stamm (Ed.), Springer, Berlin, Heidelberg (2008). 157. A. R. Burns, J. E. Houston, R. W. Carpick, T. A. Michalske, Friction and molecular deformation in the tensile regime, Phys. Rev. Lett. 82(6), 1181 (1999). 158. V. P. Gilcreest, K. A. Dawson, A. V. Gorelov, Adsorption kinetics of NIPAM-based polymers at the air−water interface as studied by pendant drop and bubble tensiometry, Phys. Chem. B, 110, 21903–21910 (2006). 159. E. Fekete, J. Moczo, B. Pukanszky, Determination of the surface characteristics of particulate fillers by inverse gas chromatography at infinite dilution: a critical approach, J. Coll. Interf. Sci., 269(1) 143–152 (2004). 160. Y. Ikada, E. Uchida, Zeta-potential of polymer surfaces, in Encyclopedia of Surface and Colloid Science, P. Somasundaran, A. Hubbard (Eds), Taylor and Francis. CRC Press. New York (2006). 161. S. Y. Fu, C. Y. Yue, X. Hu, Y. W. Mai, Analyses of the micromechanics of stress transfer in single- and multi-fiber pull-out tests, Composite Sci. Techn., 60, 569–579 (2000). 162. D. A. Mendels, Y. Leterrier, J.-A.E. Manson, The influence of internal stresses on the Microbond Test 1, Theoretical Analysis Journal of Composite Materials, 36(3), 347–363 (2002). 163. J. George, M. S. Sreekala, S. Thomas, A review on interface modification and characterization of natural fiber reinforced plastic composites, characterization methods, 6.7. Micromechanical Studies, Polymer Engineering and Science 41, 1471–1485 (2001). 164. T. Czig´any, L. M. Vas, Strength modeling of two-component hybrid fiber composites in case of simultaneous fiber failures, J. Comp. Mater., 40(19), 1735–1762 (2006). 165. J. M. Park, P. G. Kim, J. H. Jang, Z. Wang, B. S. Hwang, K. L. DeVries, Interfacial evaluation and durability of modified jute fibers/polypropylene (PP) composites using micromechanical test and acoustic emission, Composites Part B-Engineering, 39, 1042–1061 (2008). 166. T. Czig´any, B. Morlin, Z. Mezey, Interfacial adhesion in fully and partially biodegradable polymer composites examined with microdroplet test and acoustic emission, Composite Interfaces 14:(7–9), 869–878 (2007). 167. J. L. Thomason, Investigation of composite interphase using dynamic mechanical analysis: artifacts and reality, Polymer Composites, 11, 105–113 (2004). 168. R. Kumar, W. M. Cross, L. Kjerengtroen, J. J. Kellar, Fiber bias in nanoindentation of polymer matrix composites, Composite Interfaces, 11(5–6), 431–440 (2004). 169. E. Dudognon, A. Bern`es, C. Lacabanne, TSC study of length ester side group effect on sub-vitreous relaxations in poly(n-alkyl methacrylates), Polymer 43(19), 5175–5179 (2002) 170. L. C. Sawyer, D. T. Grubb, G. F. Meyers, Polymer Microscopy, Springer, Berlin, Heidelberg (2008). 171. F. Lednick´y, E. Coufalov´a, J. Hrom´adkov´a, A. Delong, V. Kolar´ık, Low-voltage TEM imaging of polymer blends, Polymer 41(13), 4909–4914 (2000). 172. P. Echlin, Handbook of Sample Preparation for Scanning Electron Microscopy and X-Ray Microanalysis, J. Heath (Ed.), Cambridge, UK Springer (2009). 173. C. Philipona, Y. Chevolot, D. L´eonard, H. J. Mathieu, H. Sigrist, A scanning near-field optical microscope approach to biomolecule patterning, Bioconjugate Chem., 12(3), 332–336 (2001). 174. G. D. Franta, I. Ohl´ıdal, P. Klapetek, R. Nepustilov´a, S. Bajer, Characterization of polymer thin films deposited on aluminum films by the combined optical method and atomic force microscopy, Surf. Interface Anal., 38, 842–846 (2006). 175. F. Hartmut, Probing the surface Tg of monodisperse PS by local thermal analysis, Macromolecules, 38, 844–850 (2005). 176. N. K. Dutta, M. D. Tran, N. R. Choudhury, Perfluoro(methylcyclohexane) plasma polymer thin film: Growth, surface morphology, and properties investigated by scanning thermal microscopy, J. Polym. Sci., Part B: Polym. Phys. 43(11), 1392–1400 (2005). 177. A. M. Belu, D. J. Graham, D. Castner, Time-of-flight secondary ion mass spectrometry: techniques and applications for the characterization of biomaterial surfaces, Biomaterials, 24, 3635–3653 (2003). 178. M. Mohai, A. Toth, P. R. Hornsby, P. A. Cusack, M. Cross, G. Marosi, XPS analysis of zinc hydroxystannate-coated hydrated fillers, Surf. Interface Anal., 34(1), 735–739 (2002).

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179. L. L. Kesmodel, High-resolution electron energy loss spectroscopy (HREELS), in Encyclopedia of Surface and Colloid Science, P. Somasundaran; A Hubbard (Eds), Taylor and Francis. CRC Press. New York (2006). 180. M. Szekeres, A. Sz´echenyi, K. St´ep´an, T. Haraszti, I. D´ek´any, Layer-by-layer self assembly preparation of layered double hydroxide/polyelectrolyte sandwiched nanofilms by surface plasmon resonance spectroscopy, Colloid Polym Sci., 283: 937–945 (2005). 181. D. Franta, I. Ohl´ıdal, P. Klapetek, P. R. I. Cabarrocas, Complete characterization of rough polymorphous silicon films by atomic force microscopy and the combined method of spectroscopic ellipsometry and spectroscopic reflectometry, Thin Solid Films, 455–456, 399–403 (2004). 182. A. Sz´ep, G. Marosi, P. Anna, J. Poti, Analysis of multicomponent systems by Raman Microscopy, Periodica Polytechnica Chemical Engineering, 47, 92–93 (2003). 183. A. Sz´ep, G. Marosfoi, G. Bertalan, A. Anna, G. Marosi, Analysis of multicomponent polymer systems by Raman Microscopy, Macromolecular Symposia, 202, 269–280 (2003). 184. A. C. De Luca, G. Rusciano, G. Pesce, S. Caserta, S. Guido, A. Sasso, Diffusion in polymer blends by Raman Microscopy, Macromolecules, 41(15), 5512–5514 (2008). ¨ 185. M. Ohman, D. Persson, C. Leygraf, In situ ATR-FTIR studies of the aluminium/polymer interface upon exposure to water and electrolyte, Progress in organic coatings, 57(1), 78–88 (2006). 186. M. J. Duer, Solid State NMR Spectroscopy: Principles and Applications, Wiley-Blackwell, 2001. 187. K. Beshah, L. K. Molnar, Characterization of interface sStructures and morphologies of heterogeneous polymers: A solid-state 1H NMR Study, Macromolecules, 33, 1036–1042 (2000). 188. A. Wernera, P. V. Konare, D. I. Svergunb, U. Hahna, Characterization of a fluorophore binding RNA aptamer by fluorescence correlation spectroscopy and small angle X-ray scattering, Analytical Biochemistry 389, 52–62 (2009). 189. O. Leupold, G. Gr¨ubel, S. V. Roth, C. Schroer, W. Roseker, M. Sikorski, A. Robert, X-ray fluorescence correlation spectroscopy – a tool to study element-specific dynamics, J.Appl. Cryst., 40, s283–s285 (2007). 190. I. K. Robinson and J. Miao, Three-dimensional coherent X-ray diffraction microscopy, MRS Bulletin, 29(3), 177–181, (2004). 191. X. Liu, F. Zheng, A. J¨urgensen, A. Perez, V. Dieste, D. Y. Petrovykh, N. L. Abbott, F. J. Himpsel, Self-assembly of biomolecules at surfaces characterized by NEXAFS Canad. J. Chem., 85(10), 793–800 (2007). 192. A. D. Smith, P. F. Schofield, G. Cressey, B. A. Cressey, P. D. Read, The development of X-ray photo-emission electron microscopy (XPEEM) for valence-state imaging of mineral intergrowths, Mineralogical Magazine; 68(6), 859–869 (2004). 193. J. Penfold, Neutron scattering for surface characterization, Current Science, 78, 1458–1467 (2000). 194. J. L. Keddie, L. J. Norton, E. J. Kramer, E. P. Giannelis, Neutron reflectometry characterization of interface width between sol-gel titanium dioxide and silicon dioxide thin films, J.Am. Ceram. Soc., 76, 2534–2538 (2005). 195. K. S¨uvegh, A. V´ertes, T. Hyodo, Positronium as a sensitive detector of changes in molecular structure, Adv. in Mol. Struct. Res. 5, 313–357 (1999). 196. C. He, E. Hamada, T. Suzuki, H. Kobayashi, K. Kondo. V. P. Shantarovich, Y. Ito, Characterization of polymer sub-surface using slow positron beam, J. Radioanal. Nuclear Chem., 255, 431–435 (2003). 197. J. Benawra, A. M. Donald and M. Shannon, Developing dual beam methods for the study of polymers, J. Phys. Conf. Ser. 126, 1–4 (2008). 198. L. A. Giannouzzi. F, A. Stevie, in Introduction to Focused Ion Beams: Instrumentation, Theory Techniques and Practic, L.A. Giannuzzi,. F.A. Stevie (Eds) Springer. New York (2005). 199. M. Draxler, S. Zoister, F. Kastner, M. Bergsmann, P. Bauer, Characterization of ultrathin Cr layers on PET by RBS and XRF Surf. Interface Anal. 34(1), 763–766 (2002). 200. G. Dev`es, S. Roudeau, A. Carmona, S. Lavielle, K. Gionnet, G. D´el´eris, R. Ortega, Fluorine microimaging and quantification using nuclear reaction analysis: A tool for validating tissue distribution of positron emission tomography tracers Appl. Phys. Lett. 95, 23701 (2009). 201. D.J. O’Connor, B. A. Sexton, R. S. C. Smart, Surface Analysis Methods in Materials Science, 23, Springer, Berlin, Heidelberg, New York (2003).

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202. K. L. Johnson, K. Kendall, A. D. Roberts, A new method for measuring the surface energie of solids, Proc. R. Soc. London A 324, 301 (1971). 203. Kinnari J. Shelat, Naba K. Dutta, Namita R. Choudhury, Interfacial interaction and morphology of EVOH and ionomer blends by scanning thermal microscopy and its correlation with barrier characteristics, Langmuir, 24(10) 5464–5473 (2008). 204. B. Bodzay, B. B. Marosfoi, T. Igricz, K. Bocz, G. Marosi, Polymer degradation studies using laser pyrolysis-FTIR microanalysis, J. Anal. Appl. Pyrolysis 85, 1–2, 313–320 (2009). 205. D. Ciprari, K. Jacob, R. Tannenbaum, Characterization of polymer nanocomposite interphase and its impact on mechanical properties, Macromolecules, 39(19) 6565–6573 (2006). 206. S. Chinn, S. DeTeresa, A. Sawvel, A. Shields, B. Balazs, R. S. Maxwell, Chemical origins of permanent set in a peroxide cured filled silicone elastomer – tensile and 1 H NMR analysis, Polymer Degradation and Stability, 91(3), 555–564 (2006). 207. X. Li, W. Fu, Y. Wang, T. Chen, X. Liu, H. Lin, P. S. Qinghua, Solid-state NMR characterization of unsaturated polyester thermoset blends containing PEO–PPO–PEO block copolymers, Polymer, 49(12), 10, 2886–2897 (2008). 208. E. Mart´ınez, E. Engel, C. L´opez-Iglesias, C. A. Mills, J. A. Planell, J. Samitier, Focused ion beam/scanning electron microscopy characterization of cell behavior on polymer micro-/nanopatterned substrates: a study of cell-substrate interactions, Micron, 39, 111–116 (2008). 209. J. R Composto, R. M. Walters, J. Genzerc, Application of ion scattering techniques to characterize polymer surfaces and interfaces Materials Science and Engineering R. 38(3–4), 107–180 (2002).

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5 Manufacturing of Multiphase Polymeric Systems Soney C. George Department of Basic Sciences, Amal Jyothi College of Engineering, Kerala, India

Sabu Thomas Centre for Nanoscience and Nanotechnology, Mahatma Gandhi University, Kerala, India

5.1 Introduction Reliable manufacturing methods and cost effectiveness are the two important factors which decide the successful production of materials. Cost effectiveness depends on higher rate of production and reliability depends on uniform quality from part to part. The successful manufacture of any product requires control over materials. Open roll mills are still being used for the manufacture of polymer blends and composites even today. However, extruders and brabenders are replacing open mills for mixing purpose. Nowadays, fully automated extruders are available. In the case of composites, the hand layup process is a primitive one but still in use. Other processes, such as compression molding, vacuum bag molding, pultrusion, filament winding and resin transfer molding are more effectively practiced these days. There are several reports available in the manufacturing process of blends and composites [1–5]. This chapter describes the fundamental aspects of various manufacturing methods of blends, composites, gels and IPNs. As well, an overview of the current trends in the processing of various polymeric materials is also briefly narrated.

5.2 Manufacturing Techniques of Polymer Blends Blends are comprised of two or more polymers, and are of commercial interest for a variety of reasons: If the mechanical properties required for a given application can be met by blending two polymers (usually

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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an expensive engineering resin with a low-cost commodity material), then the formulation becomes less expensive. If we want to recycle polymer products, we will encounter issues related to blending, since our polymer sources vary widely in terms of purity. Some polymer combinations exhibit properties that are superior to their parent polymers. This is called synergism, and it is widely used for toughening of rigid plastics. Thermodynamically there are two classes of blends; immiscible and miscible polymer blends. In immiscible blends, the constituent polymers do not mix, but remain in separate phases, leading to the formation of a dispersion of one of the polymers in a continuous matrix of the other. Experimental evidence indicates that most polymer pairs are completely immiscible. Blends can exhibit complete immiscibility and partial immiscibility, just as in small molecule systems. In miscible blends the constituent polymers mix on a molecular level, to form a homogeneous material equivalent to a polymer–polymer solution. The physical, chemical and mechanical properties are generally a weighted average of the mixture components. Polymer blends can be manufactured by different techniques such as solution blending, latex blending, mechanical blending, mechano-chemical blending, freeze drying, etc. 5.2.1

Solution Blending

In this process, polymers are dissolved in a common solvent. Blend is produced by evaporating the solvent and precipitating the resulting polymer mixture. In solution blending, selected diluents are used to dissolve the component polymers. There are several reports in literature on the solution blending of different polymer pairs. For instance, polymer blends from natural rubber and polystyrene (NR/PS) were prepared by Asaletha et al. [6]. The blends were prepared by solution casting as well as the melt blending technique. Chloroform, benzene and carbon tetrachloride were used as casting solvents. Differences in mechanical and morphological properties were observed in each case, which, in turn depends upon the interaction of the solvent with the constituent homopolymers. Mechanical properties were found to be better in samples prepared from carbon tetrachloride and it decreased in the order CCl4 , CHCl3 and C6 H6 . Poly(p-dioxanone) (PPDO)/poly(ethylene glycol) (PEG) blend was prepared by a solvent casting method using chloroform as a co-solvent [7]. Crystallization behavior of blends was monitored using differential scanning calorimetry and significant changes were observed in the crystallization pattern of blends with 5 wt% of PEG. Kohjiya and co-workers prepared copolymer blends (i-PP/EH32 and i-PP/EH57) from Isotactic polypropylene and ethylene–1-hexene by solution blending, precipitating followed by drying and hot pressing [8]. The two blends were subjected to investigation on structure and mechanical properties under uniaxial drawing. The i-PP/EH32 and i-PP/EH57 represented the immiscible and miscible blends, respectively. The tensile stresses and strains at breaking point of i-PP/EH57 were remarkably higher than those of i-PP/EH32 at room temperature. It was observed from wide-angle X-ray diffraction (WAXD) measurement that the orientation of crystallites occurred early and then propagated gradually up to a drawing ratio of 8 because chains of EH57 copolymer were incorporated into the amorphous regions between lamellae of i-PP. The miscibility and crystallization of solution casting biodegradable poly(3-hydroxybutyrate)/ poly(ethylene succinate) (PHB/PES) blends was investigated by Sun et al. [9]. The blends showed two glass transition temperatures and a depression of melting temperature of PHB with compositions in phase diagram, which indicated that the blend was partially miscible. It was found that the PHB and PES can crystallize simultaneously or upon stepwise depending on the crystallization temperatures and compositions. The spherulite growth rate of PHB increased with increasing of PES content. A series of aqueous solutions of blends based on polyethylene oxide (PEO) and hydroxypropyl methyl cellulose (HPMC) were prepared by Lee and co-workers [10]. A detailed examination of miscibility of blends was conducted by varying temperature and composition. The Huggin plots were made to study miscibility and

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it obersevd that plots were deviating from linearity according to the degree of compatibility of PEO/HPMC blends at entire concentration range. In solution blending, very small quantities of experimental polymers can be handled easily and degradation is not a problem. There are certain limitations to this method. Not all polymers are readily soluble in common solvents. Residual solvents can influence the results of analysis. It is difficult to make thick films. Moreover, in solution casting, the removal of the diluents may lead to uncertain changes in the phase morphology, thus weakening the blend. As a result of evaporation of solvent, the concentration of polymer increases and thereby phase separation occurs. The rate of evaporation plays a dominant role on the morphology and miscibility of the solution cast blends.

5.2.2

Latex Blending

In the latex blending process, polymer latices are mixed properly in a ball mill. The mixing speed and the pH are adjusted in order to get homogeneous and non-segregated latex blends. A few examples of latex blending have been cited below. The PVAc latex and NR latex blends are prepared by mixing at ambient temperature using low speed stirring [11]. A two-step mixing process is used in order to ensure preparation of homogeneous, non-segregating latex blends. The films of these blends are prepared and the film properties are investigated by scanning electron microscopy, thermal analysis and tensile testing. SEM studies have revealed the presence of phase domain structures reminiscent of the original size and shape of the poly(vinyl acetate) emulsion droplets. Samples containing PVAc as a major phase show a complex morphology. Natural rubber (NR) latex and acrylonitrile-butadiene rubber (NBR) latex blends have been prepared by latex blending [12]. The latex blends are uniformly mixed both in the latex form and in the resulting thin film. The latex blend films were formed by spin coating technique. NR/NBR blends exhibited increases in tensile strength with increasing NR content in the blends until it reaches the maximum of pure NR. The elongation at break of vulcanized NR was depressed by the addition of NBR. Latex blends of natural rubber (NR) and poly(methyl methacrylate) (PMMA) latices containing PMMA were prepared by mixing NR ammoniated latex concentrate and monodisperse PMMA latex dispersion [13]. The clarity and tackiness of the films obtained depend on the PMMA content. The morphology of these films shows asymmetric distribution of the PMMA particles across the thickness of the film. There is preferential enrichment of PMMA particles at both the air/polymer and glass/polymer interfaces and a predominating NR interior. More PMMA particles are accumulated preferentially at the glass/polymer interface than at the air/polymer interface. The accumulation of the PMMA particles at the interface is strongly dependent on the size and concentration of the PMMA particles. Latex blending has shown to yield good film formation when hard/soft latices are used. Usually the soft latex will form a film and become the continuous phase while the hard particles will act as filler and impart mechanical properties. However, it is often difficult to obtain such films because, in many instances, the hard and soft particles are not compatible, the hard particles are not uniformly distributed, or the particle packing is not well controlled. Additionally, good control of particle packing is generally achieved only for a restricted range of hard/soft particle size ratios. Beginning with well-defined latex particle morphology, the latex film formation process can be harnessed to obtain films with precisely tailored physical properties. The key requirement is that the latex particle morphology should substantially be maintained during the particle compaction and deformation stages of film formation. Thus, latexes with composite particles have received an increasing level of attention and have found applications as impact modifiers, automotive coatings, architectural coatings, membranes, etc. Latex blending offers the possibility of finer scale dispersion than solution and melt blending.

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5.2.3

Freeze-drying

A solution of the two polymers is quenched down to a very low temperature and the solvent is frozen by the freeze-drying method. The polymers will have little chance to segregate (phase separate) and will collect randomly in regions throughout the frozen solvent. Thus the state of the dilute solution is preserved. Solvent is removed by sublimation; no changes can occur because of the solid nature of the mixture. To a large extent, therefore, the resulting blend will be independent of the solvent, if the solution is single phase before freezing and the freezing occurs rapidly [14]. Freeze drying seems to work best with solvents having high symmetry. For instance, the Poly (lactic–co–glycolic acid) and hydroxyapatite (PLGA/HA) blend scaffolds were prepared by particle leaching and freeze drying techniques [15]. DMSO solutions of both PLGA and HA are evenly mixed and the resultant product is immersed in distilled water in an ultrasonic bath to leach out the salt and then the porous scaffolds were obtained. The mechanical properties of PLGA/HA blend scaffold with 2% HA are good enough to be used in cells culture and its hydrophilic surface also provides good environment for cells adhesion and growth. The SEM observation shows that cells not only grow on the surface but also migrate into the internal of the scaffold [15].

5.2.4

Mechanical Blending

Mill mixing or melt blending of constituent polymers results in mechanical polyblends. The preparation of mechanical polyblends by mill mixing or melting is problem-free compared to blending by other methods. High shearing forces for the mechanical blending of high molecular weight elastomers necessitate the use of open roll mills, internal mixers, etc. Comparable polymer viscosities at the mixing temperature are desirable for the ease of dispersion in open roll mills. The mechano-chemical blending technique is more widely used for elastomer–plastic blends and plastic–plastic blends.

5.2.4.1

Two-roll Mixing

The blending of rubbery polymers is usually done by the two-roll mixing method. In this process, one of the rubbery polymers is passed between the rollers and masticated and the second component in different percentage is added and masticated. Subsequently the additives are added in order and masticated well for the sufficient mixing time. Mixing can be done by the masterbatch technique in which the vulcanizing or curing ingredients are withheld during the mixing process. The masterbatch is converted into the compound by adding the curing agents. Roll mills are completely open to air and dust, a disadvantage, but they are the easiest mixing devices to clean. The mixing effectiveness of a two-roll mixing mill can vary from good to very poor, depending upon the rheology of the components and the skill of the operator. The blends of ethylene propylene diene monomer (EPDM) and styrene-butadiene rubber(SBR), with different crosslinking systems and fillers, are reported to be made on a two-roll mixing mill (diameter 150 mm; speed of the slow roll 24 rev/min; gear ratio 1:16). The mastication of each polymer, and the subsequent blending and compounding were done between the hot rolls of the mill at 80o C [16]. The mechanical properties of EPDM/SBR blends have been investigated with special reference to the effects of blend ratio and crosslinking systems. Among the blends, the one with 80/20 EPDM/SBR has been found to exhibit the highest tensile, tear and abrasion properties at ambient temperature. The observed changes in the mechanical properties of the blends have been correlated with the phase morphology, as attested by scanning electron micrographs. Styrene-butadiene rubber/natural rubber (SBR/NR) blends were prepared on a two roll (150 × 300 mm) at a tight nip and a friction ratio of 1:1.4 [17]. The nip gap, mill roll speed ratio and time of mixing and

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Fine morphology with compatibiliser Pellet of the secondary component

Hole formation due to interfacial instability

Coalescence

V Film of the secondary component

Rayleigh disturbances

Figure 5.1 Melt blending process in which initial morphology development in polymer blends. Reprinted from [19]. Copyright (1991) with permission from Springer Science + Business Media.

temperature of the rolls are kept constant for all mixes. The SEM studies have established the heterophase nature of the blend. SBR /NR blend with 50/50 and 30/70 blends show a co-continuous morphology [17] Compounding of natural rubber, epoxidized natural rubber-thioglycollic acid (ENR-TGA) and their blends were carried out on a two-roll mixing mill for a period of 10 min at room temperature by Okwu and Okieiman [18]. Blending natural rubber with the modified ENR with 10 parts of ENR-TGA increased the modulus by 46%. Hardness and resistance of the blend increased whereas the tensile strength and elongation at break decreased for the blend with 10 parts of ENR-TGA. [18] 5.2.4.2

Melt Blending

Melt mixing of polymeric components is usually done on internal mixers or melt mixers. High-shear mixers can generate fine dispersions, with droplet diameters <1 mm. Distributive mixing can be achieved by providing convoluted flow paths that split and reorient the flow repeatedly whereas dispersive mixing can be achieved by passing the mixture through small regions of intense deformation (Figure 5.1) [19]. Melt mixing is a complex process involving melting of the (solid) pellets, distributive mixing (particle elongation), dispersive mixing and droplet coalescence. The control of phase morphology is the key issue when desired properties are imparted to polymer blends. The shape, size and spatial distribution of the phases result from a complex interplay between viscosity (and elasticity) of the phases, interfacial properties, blend composition and processing parameters. Scott and Macosko [20], showed how and to what extent blend morphology has developed during short reactive processing by using polystyrene (matrix)-amorphous nylon (dispersed phase) blends. Morphology development at short mixing times can be summarized as follows:

r r r r

The dispersed phase forms sheets or ribbons in the matrix; Holes appear in these sheets or ribbons as a result of interfacial instability; The size and number of holes increase which leads to lace structure; Lace breaks down into irregularly shaped pieces of diameter close to the ultimate particle size.

Break-up of drops and cylinders leads to the final spherical droplets. At intermediate mixing times, thus in the melt state, mixing affects only the size of the largest particles which are transformed into smallest ones, so leading rapidly to an almost invariant morphology governed by the dynamic equilibrium. Favis studied the influence of the mixing time on the size of the dispersed phase, using a Brabender mixer [21]. It was shown for highly immiscible polypropylene/polycarbonate (PP/PC) blends that the most significant particle

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size deformation and disintegration took place within the first 2 minutes of the mixing. After 2 minutes and up to 20 minutes mixing time, very little reduction in the size of the dispersed phase was observed [22]. Both batch and continuous mixers can be used for the melt mixing process.

5.2.4.3

Batch Mixers

Common batch type internal mixers are Brabender or Haake and banbury or Moriyama dispersion mixers. Batch mixing requires only lower investment cost, but is more labor intensive in addition to low output and poor batch–batch reproducibility. It also allows the close monitoring of formulations that can combine expensive ingredients produced in small ingredients. These are also used to evaluate the compounds on a small scale. One of the drawbacks of these internal mixers is the heat conduction by the mixing shafts. A brief overview of the different parts and processing systems of DSM Micro Extruder & Injection Molder is shown in Figure 5.2. The system comprises four primary components: the micro-extruder, the feeder, miniinjection molder and the transfer cylinder. This system allows the processing of small amounts of polymers and composites into coupons suitable for physical and mechanical testing. Blends of polypropylene (PP) and ethylene-propylene rubber (EPR) and blends of polystyrene (PS) and styrene-butadiene rubber (SBR) were prepared in a laboratory-scale internal mixer at various blend compositions and rotor speed [23]. Blend morphology is studied by means of electron microscopy. For each blend pair, under the given processing conditions, the phase inversion process occurs progressively with respect to the variation in blend composition; it is within this composition range of phase inversion that dual-phase continuity is observed. In addition, characteristic torque values of blends are found to deviate negatively from a linear additivity rule; the composition range of maximum deviation from linear additivity corresponds approximately to the composition range where dual-phase continuity was observed. The effect of adding new phosphazene compounds to poly(butylene terephthalate/ polyamide 6 (PBT/PA6) blends has been studied by Bertani and co-workers [24]. The blends were prepared in a brabender batch mixer at 24◦ C at a rotor speed of 64 rpm. in order to understand the processing properties. The results indicate that there is chain extension, due to homo-coupling reaction, of PA6 when both additives are used and of PBT when 2,2-bis-(2-methoxy-4-methyleneoxy-phenoxy)-4,4,6,6,-bis[spiro(2#,2$dioxy1#,1$biphenyl)]cyclophosphazene (CP-2EPOX) is used. The intrinsic viscosity tests performed on the polymers separately added with the phosphazene compounds confirm this hypothesis. On the other side, a slight

Figure 5.2 (a) DSM mini extruder, (b) extruder open.

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compatibilising effect can be invoked for the CP-2EPOX additive due to the probable in situ formation of PBT/PA6 copolymers that act as compatibilisers for the blend.

5.2.4.4

Continuous Mixers

Continuous loading and unloading of components is the main process in continuous mixing. The advantages of continuous mixing compared to batch process are: lower power consumption; a stable process; high production volume; systematic imposition of stresses either in shear or in the shear; and elongation mode of deformation. The feed rate, screw speed, temperature as well as the discharge orifice setting and temperature can be controlled properly. The high capital investment and complex mixing are two important drawbacks of continuous mixing with high out put. Mari´c and Macosko have designed a simple cup and rotor mini-mixer, to blend very small polymer batches (0.3 g, MiniMAX) and was compared with larger lab scale extruder (5 g, DACA) and a 16 mm co-rotating twin screw (300 g/hr. PRISM) [25]. All were compared at the maximum shear rate in the cup and rotor mixer, 110 s−1 . Particle sizes of poly(propylene) [PP] dispersed in poly(styrene) (80 wt% PS) were measured by dissolving the PS, filtering and using scanning electron microscopy. The 16 mm twin screw gave somewhat smaller particle sizes than the lab scale mixers (102 μ vs 1.7 and 1.9 μ), but dispersion in the cup and rotor mini-mixer was much poorer.

5.2.4.5

Extruders

Single Screw Extruders (SSE): A single screw extruder consists of a motor drive, gear pin and a screw and the fluid layers between the screw flights and the barrel wall maintain the screw balanced and centered. A schematic sketch of single screw extruder is shown in Figure 5.3 [26]. Modern units are equipped with

Figure 5.3 A schematic sketch of extrusion process. Adapted from [26]. Copyright (2004) School of Pharmacy Shaheed Beheshti University of Medical Sciences.

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continuously variable speeds and electrically heated barrels. The different regions in SSE are solid conveying, melting and melt conveying. The melting zone is important for generation of blends with morphology, since the variation of temperature and high stress encounter within zone may profoundly affect the extrudate quality. In a single-screw extruder, the mixture is generally present in the form of striations and the striation thickness reduction quantifies the degree of mixing. In the case of preparation of blends in a screw extruder, the blending begins when the components start melting at the feed zone. At the beginning stages of mixing, the scale of segregation in the mixture is given by effective particle size of the minor component, and the mixing proceeds according to the principles of distributive mixing. In the final stages of mixing, during the passage of the polymer through the metering zone, the liquid layers/threads become unstable and disintegrate into lines of droplets caused by interfacial tension-driven hydrodynamic stability. Domain size is generally governed by a competition between droplet breakup and coalescence during melt blending. This competition depends on the viscosity ratio of blend components, processing conditions and level of compatibilization. Steady state droplet size develops during shearing due to the balance between the hydrodynamic stress which acts to deform the droplet and the restoring stress which acts to restore the spherical shape Polypropylene, and ethylene vinyl acetate copolymer were mixed on a single screw extruder and blends of polypropylene (PP) and ethylene vinyl acetate (EVA) were prepared at a processing temperature of 200o C and a screw speed of 20rpm [27]. From the visual inspection of the screw samples collected from the screw pullout experiment it was observed that the melting starts around the 6th/7th turn along the screw length. In order to characterize the morphology development prior to the breakup of the lamellar structures, the samples located at 8th, 10th and 12th channel in the feed zone were selected. The breakup of striations to droplets occurs as a result of step-up in the applied shear rate leading to application of both shear and stretching flow. There was always a sharp decrease in the number average size as compared to the volume average size for mixing in the melting stage (striation) as well as in the metering stage (droplets), showing the importance of both mean and standard deviations in the proper description of the evolution of morphology. The melt blending of the PP and scrap rubber tires (SRT) were performed in a Wortex single screw extruder (length/diameter 1 /4 32) [28]. SRT and the pellets of PP were mixed in different proportions before being added to the extruder hopper. The extruder was operated at 30 rpm with a sequence of zones temperatures at 190/200/200/210/220◦ C from feeder to die. The extruded filament was immediately quenched in water and later chopped into small granules. The unblended PP and PP/EPDM blends with or without SRT were also subjected to the same extrusion process in order to obtain a similar thermal history. When ground SRT is incorporated in PP, there is no significant change in the impact strength because of poor adhesion and large rubber particle size. However, if EPDM is present in small quantities, it acts as emulsifier at the surface of SRT particles and the two additives have a tendency to form composite EPDM-SRT particles in the PP matrix and, hence, there is an increase in the impact strength equivalent to PP/EPDM blends. Therefore, to achieve the objective of a good toughening effect, it was necessary to improve both the uniformity of rubber particles in the PP matrix and the interfacial adhesion and morphological structure between the matrix and filler, as well as forming suitable particle size and distribution. For thermal properties, in general, SRT and EPDM show the same behavior when mixed with PP–crystallinity, heat of crystallization and melting temperature were reduced. In PP/SRT/EPDM compositions, the same or an intermediate effect was observed according to SRT/EPDM proportions. Polyoxymethylene/polypropylene (POM/PP) blends of different compositions were extruded through a round die of 2 mm diameter using a single screw extruder [29]. The morphology of the extrudates was studied using a polarized light microscope (PLM), a scanning electron microscope (SEM) and the solvent extraction technique. A cable-like structure was observed in the extrudates, consisting of circular cords surrounded by a skin layer. The number and distribution of the cords matched with the holes on the breaker plate just before the die. The compositon and morphology of the cords were found to vary with the POM/PP composition ratio.

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For composition ratios between 10/90 and 40/60, the outer cords were rich in PP and only a small amount of POM fibers; particles or thin films were found. In contrast, the inner cords contained more POM fibers and they were distributed more evenly within the PP matrix. In the 50/50 blend, both POM and PP fibers were present. In blends with more than 50% POM, the cords were made up of PP fibers in a continuous POM matrix. Twin-screw Extruders (TSE): The important types of twin-screw extruders are co-rotating and counter corotating. The most commonly used is co-rotating. In the co-rotating type, the movement of the intermeshing surfaces is in opposite directions and the melt free surface is continuously renewed and the screws clean each other. This, in turn, permits to use higher screw speeds and longer barrels. Co-rotating (COR) is efficient in alternating the direction of applied stresses, providing distributive mixing by lamellae folding as well as controllable shearing (dispersing mixing). Polymer blends are melt blended mainly in TSE. This is because TSE is vastly used for reactive blending. Since SSE is not well-suited for conducting chemical reactions, the dominance is for TSEs. The SSEs are mainly used for compounding, with additional mixing units. The morphology development of the dispersed phase in PC/PE blends during the processing in a twin-screw extruder was examined by Yang and co-workers [30]. Morphology of the specimens at different positions along the screw axis was investigated by a scanning electron microscope. The effects of the extrusion temperature, viscosity ratio and the screw configuration on the development of the morphology were discussed. After analyzing the morphology development, two models of morphology development were proposed to explain the experimental results. Blends of nylon with ethylene-propylene rubber were prepared by melt mixing nylon with EPR in a corotating twin screw extruder under nitrogen atmosphere [31]. The rotor speed was 100 rev min−1 and the blending was carried out for 10 min at a temperature of 25◦ C. The liquid permeability of the blends is found to be affected by the morphology of the blends. In EPR/nylon 30/70 blend, EPR was dispersed as domains in the continuous nylon matrix. The continuous nylon matrix restricts the swelling and hence the solvent uptake is negligibily small. But in EPR/nylon 50/50, an interpenetrating morphology is observed. Here both EPR and nylon form a co-continous morphology. In EPR/nylon 70/30, nylon is dispersed as domians in the continuous EPR matrix. The morphology of EPR/nylon 50/50 and EPR/nylon 70/30 system significantly accelerates the swelling process. The evolution of phase morphology in uncompatibilized and reactively compatibilized nylon-6/ethylene propylene rubber (EPR) blends has been investigated by scanning electron microscopy with special reference to the effect of mixer type and size [32]. Three types of mixing instruments were used for the melt preparation of the blends. These include the Haake Rheocord mixer, the DSM twin-screw mini extruder and the industrial Werner Pfleiderer twin-screw extruder (ZSK-25). In the reactively compatibilized blends, EPR-g-maleic anhydride (MA) has been used as compatibilizer precursor. The MA group of the EPR reacts with the amino group of nylon, forming a graft copolymer at the interface, which decreases the interfacial tension and suppresses the coalescence. In the case of uncompatibilized blends the final morphology is developed within 3–4 min of the mixing time in the DSM twin-screw mini extruder. The size of the dispersed phase in the uncompatibilized blends, as obtained in the Werner Pfleiderer twin-screw extruder and the DSM mini extruder, was found to be smaller than that obtained in the Haake Rheocord mixer on account of elongational flow occurring in the twin-screw extruder. In the case of the Werner Pfleiderer twin-screw extruder, sampling was done at various points in order to understand the evolution of the phase morphology along the axis of the extruder. It was found that the final morphology of the uncompatibilized blends in the Werner Pfleiderer extruder was controlled by the geometry of the die at the exit. In the case of reactively compatibilized blends, the morphology development in the DSM twin-screw mini extruder was so fast that the final morphology developed within 30 sec of mixing. This occurred because since the compatibilizer formed as a result of the

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reaction remained located at the blend interface, coalescence was substantially suppressed, and thereby the blend system rapidly attained a stable morphology. The evolution of blend morphology during compounding in a twin-screw extruder was investigated for blend systems such as (i) polystyrene (PS)/poly(methyl methacrylate) (PMMA), (ii) PS/polycarbonate (PC), (iii) PS/high-density polyethylene (HDPE), and (iv) PS/polypropylene (PP), by studying the effects of viscosity ratio, blend composition, and processing variables [33]. The viscosities of the five polymers (PS, PMMA, PC, HDPE, and PP) chosen for melt blending were measured over a wide range of temperatures at shear rates ranging from 0.001 to 1000 s−1 . A ‘screw pullout’ experiment was conducted to investigate the evolution of blend morphology, determined by transmission or scanning electron microscopy, along the extruder axis. They found that the initial blend morphology depends very much on the difference in critical flow temperature (Tcf) or melting temperature (Tm) between the constituent components and, also, on the viscosities of the constituent components. They observed that a co-continuous morphology was formed at the front end of the extruder, which then transformed into a dispersed morphology towards the end of the extruder. They found that the blend ratio determined the state of dispersion for asymmetric blend compositions and the viscosity ratio determined the state of dispersion for the symmetric blend composition. Non reactive and reactive polyamide /polysulfone (PA/ PSU) blends prepared in a twin- screw mini extruder and Haake batch mixer [34]. The study indicates that at the given processing condition the miniscrew extruder is more effective in mixing both non-reactive and reactive PA/PSU blends than the Haake batch mixer. Therefore batch mixer is a good screening device to prepare blends of newly developed polymers which have high melt viscosity and are available in rather limited amount. 5.2.5

Mechano-chemical Blending

In this method of blending, two polymers are mixed by melt mixing and then crosslinked using a cocrosslinker. This intercrosslinking between the two polymers improves the phase morphology by the proper control of domain size and distribution thereby providing resistance to phase separation. Ning-Jian and Qi have prepared polyamide-6/polypropylene/styrene-butadiene styrene (PA-6/PP/SBS) blends by solid-state mechanochemical method [35]. In this process, the solid was pulverized at ambient temperature under the strong pressure and shearing stress in pan milling type equipment. It was attempted to control the phase structures of the plastics through compatibilization by milling and changing processing temperatures, so as to improve the phase interface and properties of the PA/PP/SBS blend. It was found that pan milling can improve the compatibility of PA with SBS and PP while pulverizes PA6/PP particles into ultrafine powder. While blending the PA/PP ultrafine powder with SBS a low loading of PA/PP contents is required to increase the tensile strength of the matrix from 12.7 MPa to 25.0 MPa. Recently Prut has reviewed the chemical processes that occur during the processing of polymer materials under the action of mechanical forces [36]. The chemical modification and blending of polymers in an extruder reactor were considered. 5.2.6

Manufacturing of Polymer Blends Using Supercritical Fluids

Supercritical carbon dioxide is found to be used for blending and processing of polymers. It lowers the melt viscosities of polymer melts and allow more effective blending of polymers. This phenomenon occurs through a free volume mechanism. Carbon dioxide will absorb between the polymer chains causing an increase in free volume, resulting in a decrease in chain entanglement. The increase in free volume will also increase chain mobility, which ultimately reduces melt viscosity [37]. The objective of using supercritical CO2 in polymer blending is to have a larger amount of carbon dioxide in the high viscosity material and lower its viscosity more than the other material. One way to accomplish this is for the high viscosity material to have a higher affinity for carbon dioxide than the material with the lower viscosity. This will cause the viscosity ratio of

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the polymer melts to become closer to 1. A finer dispersion of the minor component may result. A reduction in the interfacial surface tension of the polymers caused by the addition of supercritical carbon dioxide may also contribute to a finer dispersion of the minor component. One of the potential applications of supercritical carbon dioxide is to manipulate the viscosity ratio of polymer blends. An increase in impact strength may be obtained by finely dispersing the minor phase in the major component. Expanding the spectrum of melt viscosities in polymer blending can lead to new and improved polymer blends. Altering the final morphology of blends by using this technique can be favorable in applications where properties such as permeability, strength, and ease of processing depend strongly on blend morphology. Blends of poly(acrylic acid)/nylon6 and polystyrene/nylon6 were prepared using supercritical CO2 as substrate-swelling agent and monomer/initiator carrier [38]. Both supercritical CO2 /nylon6 binary system and SC CO2 /monomer/nylon6 ternary system were studied in detail using differential scanning calorimetry, infrared spectroscopy, and polarizing microscopy. Supercritical CO2 -induced crystallization was observed in modified nylon6. The preparation of polysulfone/polycaprolactone (PS/PCL) blend membranes using a CO2 -assisted phase inversion method was reported [39]. Membranes containing from 0 to 50 wt% of PCL were prepared and analysed in terms of morphology, miscibility, hydrophilicity, transport properties and mechanical performance. Membrane porous structures were characterised by scanning electron microscopy (SEM), showing that significant changes in the morphological characteristics were obtained upon the addition of PCL. The water flux measurements confirmed higher porosity and permeability for membranes with higher PCL weight ratio but contact angles indicated that the hydrophobic nature of the surfaces was increased. Differential scanning calorimetry (DSC) and dynamic mechanical analysis (DMA) experiments showed that PCL addition increased the pores non-homogeneity, improved the damping properties of the membranes and decreased the elastic behavior. Polymer blends have been prepared in a nonreactive way by batch mixing and extrusion with supercritical CO2 as a plasticizing agent [40]. The dissolution of supercritical CO2 in blends leads to decreased shear thinning and a finer dispersion of the minor component. For example, effect of supercritical CO2 on the dispersion of the minor phase, PMMA, in a PMMA/PS blend [41] has been reported. The greater reduction in viscosity of the minor component allows better momentum transfer from the more viscous major component and causes the minor component to break up into smaller droplets. The effect of CO2 on a dispersed phase can be seen through earlier completion of phase inversion compared to blending without CO2 [42]. Phase inversion is a transformation of high-melting polymer from a disperse phase to a continuous phase. In the presence of CO2 , the length of the extruder required for the phase inversion is shortened because of the decrease in glass transition temperature of the high-melting polymer, and the rest of the extruder length is effectively utilized for reduction of the droplet size of the disperse phase. These studies clearly confirm the possibility of controlling blend morphology (and, therefore, the properties of the blend) in the presence of supercritical CO2 . If foaming of a polymer blend is not desired, venting of CO2 from the blend is necessary which may cause demixing of the polymers To prevent demixing, additives such as carbon black, calcium carbonate and nano-clay particles can be used.

5.3 Manufacturing Techniques of Polymer Composites As an alternative for metallic materials, lightweight but extremely tough and rigid composites can be prepared by the combination of polymers and fibers or fillers. Based on the demand of the materials, it is possible to choose the most attractive and most efficient manufacturing processes. All important manufacturing process are described here.

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5.3.1

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Hand Layup Process

Hand layup, or contact molding, is the effective though simple process on which the composites industry was founded nearly 60 years ago. It continues to be very extensively employed. In this molding process, a release agent is applied on the mold. It is followed by a resin-rich pigmented gel coat and when this is cured a thin layer of matrix resin is applied to the tool surface prior to the positioning of the fiber reinforcement. The fiber reinforcement usually is chopped strand mat (CSM), although the entire or partial use of fabrics such as woven roving can apply. The layers of reinforcement and the resin matrix control the thickness of the laminate. The laminate is consolidated using either a brush or a roller. Just before full-cure has been developed, the component may be roughly trimmed to shape, after which it is stripped (or released) from the mold-tool, which is then cleaned, released and returned to use. The two most common resins used in hand layup are polyesters and epoxies. Polyesters are favored because of low cost, wide availability, and ease of handling. Also, polyesters can be easily removed from a wide range of molds with standard mold release agent. Polyesters usually contain fillers such as calcium carbonate or talc to prevent large volumetric shrinkage on curing, which can lead to large warpage, distortion, or ripple surfaces. Epoxies are more expensive but exhibit better adhesion to reinforcements. Mold release agents are needed in the formulation otherwise the parts would stick to the mold, leading to difficulties of demolding, distortion, warpage or even severe damage of the molded parts. Both matrix materials can be formulated to be flame retardant, non-burning, or self-extinguishing. Epoxy parts are used mainly for dimensionally stable, high-strength applications, while polyesters are used in high-volume, less critical applications. Glass fiber is the most common reinforcing material for hand layup of polyester and epoxy matrix composites. The fibers are made up of bundles of fine filaments, which have been pretreated with silane couple agents in an attempt to achieve good wetting and interactions. In hand layup, the most common glass fiber forms are roving, woven roving, cloth and mat, chopped and milled fibers, and veils. The fiber form is not so critical as in other processes such as compression molding or filament winding. The major advantages of using unsaturated polyester and epoxy systems with hand layup techniques are: ease of fabrication; low-cost tooling; wide range of available colours; and good composite properties and large size capability Thomas et al. [43, 44] have prepared several thermoset composites by the hand layup process and cured by compression molding technique. One among them is an oil palm fiber reinforced PF composite [43]. It was prepared by a prepreg route. The fibers were chopped into desired lengths (10, 20, 30 and 40 mm) and after being dried in an oven were spread in a mold cavity. Mold was closed and pressure applied to form a single unit. The mat was then impregnated in the resin and the prepreg was kept at room temperature up to a semicured stage. It was then pressed at 100◦ C to get a three dimensionally crosslinked network. The main advantages of the hand layup process are low capital outlay, secondary bonding, no size limit and flexibility, whereas the limitations are that the process is operator-dependent, labour-intensive, low production rate, poor weight and thickness control, in addition to the fact that only one molded face is possible. Hand layup molding is used for the production of parts of any dimensions, for example, technical parts of a few tens of m2 area, as well as swimming pools of 150 m2 developed area. It is recommended for small and medium volumes, for which the investment in molds and equipment is not very high. These include larger structural parts such as: boats, portable toilets, picnic tables, car bodies, diesel truck cabs, hard shell truck bed covers and high performance air craft skins and interiors. The hand lay up process is labor-intensive plus the plastic resins produce toxic fumes requiring well-ventilated facilities and protective equipment for workers.

5.3.2

Spray Layup Process

In spray layup process, both resin and fibers are sprayed on to the mold using a spray gun instead of by a hand-in-hand layup process. Chopped fiber and resin are placed appropriately over the mold. The main

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advantages are: low material cost; high production rate; low tooling cost; and large parts. The limitations are: very operator-dependent, very poor thickness control, single molded face and random reinforcement. Resins used in spray up are usually unsaturated polyesters or epoxies. An excellent surface finish can be achieved by spraying gel coat onto the mold. Woven roving can be added during the sprayup process for strength improvement in a specific direction or location. Materials for spray up molds may be wood, vinyl polyester, epoxy, rubber, or steel. The sprayup process is ideal for low- to medium-volume applications and is particularly well suited for efficient fabrication of large shapes. Additionally, the decision to choose sprayup over hand layup is based on both the mechanical properties required and the processing constraints. Parts produced by sprayup are boat hulls, swimming pools, covers, end caps of large diameter tanks, reinforced acrylic bathtubs, shower units, recreational vehicle components, truck roofs, large panels, containers, housings and guards, furniture, and simulated masonry. Sprayup systems are often used to back up thermoformed acrylic plastic sheet for bathtubs, shower stalls, hot tubs, camper tops, cycle fairing, and pool accessories, etc. 5.3.3

Vacuum Bag Molding

Vacuum bag techniques have been developed mainly for fabricating complex shapes, double contours and relatively large components. A flexible sealed bag under which a vacuum is drawn and exposes the laminate in the mold below to atmospheric pressure. The assembly is then heated in an oven to promote flow and cure the resin. Low-cost equipment and tooling is one of the attractive features of this process. A schematic sketch of vacuum bag process is shown in Figure 5.4 [45]. Vacuum bagging uses atmospheric pressure to press the cloth tightly against the surface being covered so that the excess resin is squeezed out and soaked up in a disposable outer wrap. This technique requires a vacuum bag, a vacuum pump capable of pulling a significant vacuum and various accessories and supplies. Thus vacuum bagging has been mostly restricted to large-scale commercial use. The traditional single vacuum bag (SVB) process is best suited for molding epoxy matrix-based composites because of their superior flow and the absence of reaction by products or other volatiles [46]. This is not the case for other classes of materials such as polyimides and phenolics. Polyimides and phenolics are cured by condensation reactions which generate water as a reaction byproduct. In addition, these materials are commonly synthesized as oligomers using solvents to facilitate processability. SVB molding, without additional pressure, normally fails to yield void-free quality composites for these classes of resin systems. A double vacuum bag (DVB) process using common molding equipment was designed for volatile management in composite fabrication. This experimental DVB process affords superior volatiles management compared with the traditional SVB process. Void-free composites are consistently fabricated as measured by C-scan

Figure 5.4 A schematic sketch of vacuum bag process. Reproduced with permission from www.azom.com, Article 401, http://www.azom.com/article.aspx?ArticleID=401 Copyright (2011) Azom.com.

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and optical photomicroscopy for high-performance polyimide and phenolic resins. Double bagging provides vacuum integrity and controls bag relaxation while flow media controls the flow front to allow high quality aerospace-grade products The multi-layer woven silk/epoxy composites were produced by a vacuum bagging process in an autoclave with increasing layers of silk fiber [47]. Two sets of sample were prepared. In the first set, silk was treated with a surface treatment using a silane-based coupling agent and in the second set, the silk fiber was not treated at all. All the samples were tested for interlaminar fracture using the double cantilever beam specimens testing method. Stable crack propagation was observed during the tests and the crack propagation areas for the untreated fibers showed that all the fibers were bare with no matrix as observed in scanning electron microscopy. Failure occurred at the fiber matrix interface with no fibere bridging observed between the two fracture surfaces. M˚anson and co-workers have investigated the effect of honeycomb pressure level on skin–core adhesion and skin quality during the vacuum bag processing of honeycomb structures [48]. The pressure level inside the honeycomb cells plays an important role and is controlled by the permeability to air of the skins. An initial range of skin permeability to air was explored by perforating the prepregs and adhesive layer selectively. The role of the resulting pressure inside the honeycomb on skin–core adhesion and skin quality was evaluated. Prepreg air permeability was found to control skin–core adhesion through the pressure in the honeycomb and potential outgassing of the adhesive layer. An optimal range of initial pressure inside the honeycomb was found to be 40–70 kPa. A universal process window was proposed to determine the time frame of vacuum application leading to an optimal initial honeycomb pressure level. The main advantages of this process are: low capital outlay; low cost tooling; and usage of large components well suited to making sandwich panels. That the process is labor intensive with low production rate and only one accurate surface is obtained are cited as disadvantages. 5.3.4

Resin Transfer Molding

Resin transfer molding (RTM) is a low-pressure closed molding process, where a mixed resin and catalyst are injected into a closed mold containing a fiber pack or perform. After the resin is cured, the mold can be opened and the finished component removed (Figure 5.5) [49]. A wide range of resin systems can be used including polyester, vinylester, epoxy, phenolic and methyl methacrylates, combined with pigments and fillers including aluminum trihydrates and calcium carbonates, if required. The fiber pack can be glass, carbon, natural fibers, aramid, or a combination of these. The closed molding process is cleaner and healthier and attracts higher skilled employees. The dynamic mechanical properties of treated sisal fiber-reinforced polyester composites fabricated by resin transfer molding (RTM) have been studied with reference to fiber surface modifications, frequency and temperature [50]. The sisal fibers have been subjected to various chemical and physical treatments like mercerization, heating at 100◦ C, permanganate, benzoylation and vinyl tris(2-ethoxymethoxy) silane to improve the interfacial bonding with isophthalic polyester resin. Results indicated that treatment changed the storage modulus (E ), loss modulus (E ) and damping factor (tan δ) drastically at a wide range of temperature. Resin infusion was modeled and analytic solutions were obtained for vacuum assisted resin transfer molding (VARTM) [51]. Compaction behavior of the fiber preform was examined experimentally and the influence of compressibility of the preform on the resin infusion was investigated mathematically. The analytic model proposed in the present study predicted flow front advancement through the preform. The model provided pressure and thickness distributions of the region impregnated by the resin. Tari et al. have devised a methodology for the rapid process development, resin transfer molding (RTM) in order to reduce product development cycle time and lower the cost of developing composite parts [52]. As a first step towards this goal the feasibility of an alternative rapid prototyping technique, namely laminated

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Figure 5.5 A schematic sketch of Resin transfer process. Reprinted from http://www.ctihuatai.com/. Copyright CTiHuatai 2008.

object manufacturing (LOM), to fabricate molds has been investigated. The durability and dimensional stability of such a mold, as well as its economic practicality in the RTM manufacturing of a composite I-beam were considered. To keep tooling and fixturing simple, vacuum pressure was used to draw’resin into the mold. However, injection at this low pressure proved difficult. To remedy the situation a high permeability layer (HPL) was incorporated into the process. The benefits and drawbacks of utilizing such a layer are also discussed in the paper. Furthermore, an RTM simulation program developed at UCLA was employed to determine optimum gate locations. Finally, the total hours to manufacture the first acceptable part were quantified for comparison with future studies. The current study has shown that, when properly treated, LOM molds can be used in conjunction with the RTM process to produce prototype composite parts. Since the first mold showed no signs of delamination or cracking after the fabrication of the three I-beams, it can be concluded that such a mold is durable enough for prototyping purposes. The effects of glass fiber surface modification on the flow characteristics of unsaturated polyester (UPE) resin were investigated in the resin-transfer molding (RTM) process [53]. g-Methacryloxypropyl trimethoxy silane (g-MPS) was used as a glass fiber surface modifier.It was found that surface energy of glass fiber was decreased by g-MPS treatment by advancing contact-angle measurement. Unsteady state permeability of glass fabric preforms was measured according to Darcy’s law. The apparent permeability of g-MPS-treated glass fabric preforms was slightly lower than that of untreated fabric performs because of the macro/micro flow induced by capillary action. The void contents and the flexural properties of the cured glass-fiber/UPE composites were estimated and morphological study of the glass-fiber/UPE composites was also performed by SEM. When the fiber surface was treated with g-MPS, the void content and the flexural properties of the glass-fiber/UPE composites were different in different regions of the mold cavity. The present series of investigations has been undertaken to characterize various effects of surface treatment on liquid resin flow in the resin transfer molding process. From the study of flow characteristics of unsaturated polyester (UPE) resin in glass fabric preform, the following conclusions were obtained. RTM has a very low tooling cost and simple mold clamping requirements. RTM has its ability to encapsulate metal ribs, stiffeners, and inserts and so on within the molded laminate. RTM process has been used in molding parts such as cabinet walls, chair or bench seats, hoppers, water tanks, bath tubs and boat hulls. RTM has advantages of producing low-cost composite parts with complex and large structures in near net shapes. Relatively fast cycle times with good surface definition and appearance can easily be achieved. The capability

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to consolidate parts in the closed mold at elevated temperatures allows the saving of considerable amount of time and trouble over conventional lay up processes. Since RTM is not limited by the size and pressure associated with autoclave, new tooling approaches can be utilized to fabricate large, complicated structures. On the other hand, RTM is suffering from some disadvantages. The mold design is critical and requires good tools or great skill. Improper gating or venting may result in defects. Flow pattern and resin impregnation uniformity control is difficult. Radii and edges tend to be resin-rich. Reinforcement movement and edge flow during resin injection are sometimes a problem. 5.3.5

Pultrusion

The reinforcement is impregnated with resin and pulled through a heated die, which gives shape to the final product. It emerges from the die fully cured or by injection directly into the die. The main advantages of this process are high mechanical properties, high production rate, etc. Figure 5.6 [45] shows the details of the pultrusion process. The pultrusion of glass-reinforced polypropylene and its behavior on the final properties were studied [54]. The experiment was conducted on a new pultruding machine. The process variables investigated during the experiment were preheating temperature, heated die temperature, cooled die temperature and pulling speed. Their effect was assessed by the characterization of pultruded sections based upon fiber distribution over a cross-section, flexural modulus and strength, surface finish and pulling force. The pultrusion of a flax-reinforced polypropylene commingled yarn containing discontinuous flax and polypropylene fibers, was investigated [55]. This was the first attempt to pultrude this material. Rectangular cross-sectional profiles were successfully produced using a self-designed pultrusion line. In a series of experiments carried out with yarns of two different flax fiber contents, the pultrusion parameters were varied. The preheating and die temperatures and pulling speed are the most relevant parameters, which decide the pultrusion of natural fiber composite profiles at industrial scale. A complete characterization of each profile was conducted in order to examine the influence of processing parameters on profile quality. The high mechanical properties of the composites obtained by pultrusion result from the orientation of long fibers The fatigue properties of E-glass-type fiber-reinforced phenolic composites manufactured by the pultrusion and the hand lay up processes were compared [56]. For the hand lay up process fatigue data are shown for composites with fibers aligned in the 0◦ longitudinal tensile loading directions, 90/0/90 orientations and 0/90/0 fibers with tensile loading applied also in the longitudinal direction. In the pultruded specimens fatigue life data are presented at room temperature and only one value of stress ratio and frequency. Pultruded composites have adequate fatigue strength values but lower than the results obtained for the longitudinally loaded 0◦ hand lay up specimens. The hand lay up composite of the unidirectional fiber orientation (0◦ ) have shown higher values of fatigue strength at the elevated temperature of 200◦ C in comparison with the room temperature values.

Figure 5.6 A schematic sketch of pultrusion process. Reproduced with permission from www.azom.com, Article 401, http://www.azom.com/article.aspx?ArticleID=401 Copyright (2011) Azom.com.

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The advantages are: high mechanical properties; high production rate; close tolerance; consistency. The disadvantages are: shape limitation and limited transverse properties. Pipe, tubing, structural beams and special shapes are nice examples of pultruded products. The products are heavy with high strengths. They are often used to replace metal structural parts when a corrosive resistant or non-magnetic material is required. More than 90% of all pultruded products are glass fiber-reinforced unsaturated polyester. Other resin and matrix systems are employed when higher demanding performance is required. For example, vinyl ester resins are used for greater corrosion resistance and epoxy resins are used for a combination of superior mechanical and electrical properties. When high service temperature and high mechanical properties are required, aramid or carbon fibers with epoxy resin are used. 5.3.6

Filament Winding Process

Filament winding is a simple and effective method for producing products such as pipes and cylindrical containers in a wide range of sizes. The filament winding process was originally invented to produce missile casings, nose cones and fuselage structures; however, with the passage of time, industries other than defense and aerospace have discovered the strength and versatility of filament winding. It is a noteworthy method for the processing of reinforced plastics. Storage tanks and pipes for the chemical and other industrial areas are the main applications of the process. Polyester resin and fiberglass are the major constituent materials. The process consists of wrapping bands of continuous fiber or roving over a mandrel in a controlled operation. Many layers of the same or different patterns are placed on the mandrel, and the repetitive patterns and reinforcement spacing are subject to close control. During winding, fiber tension generates pressure between layers of uncured composite. This pressure influences the compaction and void content of the article, which in turn controls the strength and stiffness of the wound products. The resin may be added before, during, or after the winding operation. Finally the resin is cured at room or elevated temperatures without pressure and the mandrel is removed from the component after curing if the mandrel is not an element of the structure. Finishing operations such as machining or grinding are usually not necessary. The method of manufacturing of wound pipe-like structures is described in Figure 5.7 [45]. A new type of filament wound, flat strip composite of polyethylene fiber reinforced polyolefins were prepared by Kazanci et al. [57]. The versatility of this product is demonstrated by experimenting with different polyolefin compositions based on ethylene-butene copolymers, with a range of fiber volume fractions and various winding angles. The mechanical performance of the composite product agrees with standard theoretical predictions for angle-ply composites, while the dynamic mechanical analysis reflects the viscoelastic nature

Roving supply

Tensioner Resin bath Tool

Semi-finished component

Shuttle

Axle

Track

Figure 5.7 A schematic of film winding process. Reproduced with permission from www.azom.com, Article 401, http://www.azom.com/article.aspx?ArticleID=401 Copyright (2011) Azom.com.

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of the matrix. In particular, the branching density of the polyolefin backbone, expressed by the β-transition peak, dominates the dynamic mechanical behavior. All PP composites pipes or rings are manufactured in a simple filament winding process in which the co-extruded tapes (or woven fabrics) are cold-wound onto a mandrel, and subsequently heated to achieve consolidation [58]. This consolidation of tape or fabric layers occurs due to the melting of adjacent PP copolymer skin layers of neighboring tapes, which solidify on subsequent cooling to result in a coherent load-bearing structure of bonded, highly-oriented PP tapes. During consolidation of filament wound tapes, hoop tensile stresses and radial compressive stresses are generated in the tape because the tape has a negative coefficient of thermal expansion whereas the mandrel has a positive thermal expansion. These hoop tensile stresses prevent tape shrinkage while the radial compressive stresses favor consolidation of the pipes. No rotation of the mandrel is required in the oven during consolidation. The main advantages are: excellent mechanical properties; high production rate; good control of fiber orientation; good thickness control; good fiber content control; good internal finish. The disadvantages are: a limited range of shapes and a limited number of practical winding patterns. Filament wound items are produced in quantity for both aerospace applications and the commercial market. Winding technology and design are different for these two fields because of different performance requirements. Aerospace winding has become highly specialized and is designed for superior performance at a premium price. The major efforts are directed to such military applications as rocket-motor cases, high pressure vessels, missiles launch tubes, and helicopter blades. Here the strength-to-density ratio is of paramount importance, and the aim is to orient the fiber along the direction of the applied stress. The commercial sector, on the other hand, looks for lower production costs and moderate strength. An unsaturated polyester/glass fiber system is preferably employed in combination with simplified winding equipment to favor high production speeds and low cost. Corrosion resistant and electrical grade materials are used in markets for storage tanks, reinforced pipes, fire-resistant pressure vessels, chimney liners and stacks, lightning arresters, chemical storage and processing tanks, springs, drive shafts, and wind turbine blades. 5.3.7

Reaction Injection Molding

Over the last 20 years, reaction injection molding (RIM) has emerged as a popular method of producing premium quality molded polyurethane parts. These range from simple parts to complicated structural components, typically in volumes of one to 5,000 pieces. Composite systems can be foamed or solid, rigid or elastomeric, but are molded with fiber reinforcements such as glass to enhance the structural properties of the molded part. Stiffness and impact strength are enhanced by adding reinforcement in the material stream (Reinforced RIM-RRIM), or by using a molded preform in the mold that is encapsulated (Structural RIM-SRIM). Another RIM-processed material showing some promising results for specific applications is polyurea. Modified resins, particularly those modified with rubber, show good stiffness and a good improvement in impact strength. Reinforced Reaction Injection Molding or RRIM is a process used to produce polyurethane and polyurea thermoset polymers. When short fibers (1.6 mm), carbon or mineral fillers are incorporated into one of the liquids, it increases modulus and reduces coefficient of expansion. A schematic representation of the RRIM process is shown in Figure 5.8 [59]. Both the resin slurry and the isocyanate prepolymer in the tanks are usually circulated (by pump) through the RIM dispensing equipment and the mixhead before returning to the tank. This circulation, in combination with in-tank agitation, ensures homogeneity of the components. 5.3.8

Rotational Molding

Rotational molding (or rotomolding) is a method used to mold hollow parts, especially those with complex and varied shapes not easily obtainable by other hollow-art processes. It involves the slow tumbling, heating,

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Figure 5.8 A schematic sketch of reinforced reaction injection molding process. Courtesty of Bayer MaterialScience AG Group.

and melting of a thermoplastic powder in a biaxially rotating mold to produce seamless, hollow plastic parts and composites. It is also a virtually shear-free and pressure-free process. Lower viscosity resins are typically used in this process as they are more readily sintered to ensure a good surface finish. Rotational molding offers design advantages over other molding processes. With proper design, parts assembled from several pieces can be molded as one part, eliminating high fabrication costs. The ability to add prefinished pieces to the mold alone is a great advantage. Since the mold require less tooling, it can be manufactured and put into production much more quickly than with other molding processes. Rotational molding is also the desired process for short runs. The wall thickness uniformity and part weight can be easily maintained. There is very little waste of material due to scrap. As the rotomolding is a low-pressure process, sometimes designers face hard-to-reach areas in the mold. Unlike other processes where only the product needs to be cooled before being removed, with rotational molding the entire mold must be cooled. Products that can be manufactured using rotational molding include storage tanks, bins and refuse containers, airplane parts, doll parts, road cones, footballs, helmets, rowing boats and kayak hulls. Playground slides and roofs are also generally rotomolded.

5.4 Manufacturing Techniques of Nanocomposites Nanocomposites are mostly manufactured by three important methods: (a) solution intercalation; (b) in situ intercalation; and (c) melt intercalation. 5.4.1

Solution Intercalation

The polymer or pre-polymer is soluble and the nanofiller usually is a layered one, such as silicates and hydroxides which are swellable in a suitable solvent. The layered filler is first swollen in a solvent such as water, chloroform or toluene. When the polymer and layered material solution are mixed, the polymer chains intercalate and displace the solvent within the layers. By the removal of the solvent, the remained intercalated structure leads to a polymer-layered nanocomposite.

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Ethylene vinyl acetate/expanded graphite (EVA/EG) nanocomposites were prepared by the solution intercalation technique followed by compression molding, and the resulting nanocomposite properties were compared with those of natural graphite (NG) filled EVA composite [60]. TEM pictures revealed that the graphite platelets were homogeneously distributed in the EVA matrix, whereas the NG particles formed agglomerates inside the rubber matrix. With a loading of 4 wt% expanded graphite in EVA, the maximum tensile strength and the modulus improved over the neat EVA film by ∼35 and 150% respectively. But at a higher loading of 8 wt%, the tensile strength and elongation at break exhibited significant decrement, which may be due to the relative agglomeration of the graphite particles within the matrix. The incorporation of expanded graphite provided tremendous improvement in thermal conductivity and thermal degradation resistance. A maximum shift of 14◦ C in maximum rate of degradation was observed for 4 phr EG addition. Ethylene vinyl acetate and carbon nanofiber nanocomposites (CNF) were prepared by the solution intercalation technique [61]. The reinforcement of EVA matrix by CNFs was studied by analyzing various properties For instance, the tensile strength was greatly increased (61%), even for very low fiber content (i.e., 1.0 wt %). The surface modification of the fiber by high-energy electron beam and gamma irradiation led to better dispersion in the rubber matrix. This in turn gave rise to further improvements in mechanical and dynamic mechanical properties of EVA. The thermal conductivity also exhibited improvements from that of the neat elastomer, although thermal stability of the nanocomposites was not significantly altered by the functionalization of CNFs. The disadvantages of this process are that a relatively large number of solvent molecules have to be desorbed from the host to accommodate the incoming polymer chains. The desorbed solvent molecules gain one translational degree of freedom and the resulting entropy gain compensates for the decrease in conformational entropy of the confined polymer chains. The latex stage compounding technique is a commonly used method for manufacturing nanocomposites. As most rubbers exist in latex forms, latex particles with sizes of 50 nm are uniformly stabilized in water which offers a good performance/cost ratio.

5.4.2

In Situ Intercalative Polymerization Method

In this method, the layered material is swollen within the liquid monomer or a monomer solution so that the polymer formation can occur between the intercalated sheets. Polymerizations can be initiated either by heat or radiation or by the diffusion of suitable initiator or by an organic initiator or catalyst fixed through cation exchange inside the interlayer before the swelling step. Polyurethane/nanosilica composites were prepared using polyester polyol/nanosilica composite resins obtained from in situ polymerization or blending methods [62]. It was investigated by fourier transform infrared spectra (FTIR), dynamical mechanical analysis (DMA), transmission electron microscopy (TEM), contact angle measurement, X-ray photoelectron spectroscopy (XPS) and atomic force microscopy (AFM), respectively. It was found that more polyester segments had chemically bonded with silica particles during in situ polymerization than during blending. The introduction of nanosilica increased the glass transition temperatures (Tg) of polyurethanes, and different preparation methods and different particle sizes caused various impact on the Tg values. The polymethyl acrylate/europium oxide (PMA/Eu2 O3 ) porous and layered nanocomposite was prepared by in situ polymerization and characterized by means of X-ray diffraction (XRD), transmission electron microscopy (TEM), scanning electron microscopy (SEM), and infrared ray (IR) spectroscopic techniques [63]. Microscopic investigation of the nanocomposite was carried out by atomic force microscopy (AFM). The results showed that the shape of the composite was layered and porous. Eu2 O3 was grafted when methyl acrylate (MA) polymerized and Eu2 O3 particles appeared on both sides of the chains of PMA

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143

Melt Intercalation or Melt Blending Method

In this method, a mixture of polymer and nanofiller is annealed statically or under shear above the softening point of the polymer. This method has greater advantages over either in situ intercalative polymerization or the polymer solution intercalation. First, this method is environmentally benign due to the absence of organic solvents. Second, it is compatible with current industrial processes such as extrusion and injection molding. The melt intercalation method allows the use of polymers which were previously not suitable for in situ polymerization or solution intercalation. Chen and co-workers have prepared Poly-L-Lactic acid/poly caprolactone (Organomontmorillonite PLLA/PCL/OMMT) nanocomposites by blending PLLA, PCL and OMMT [64] and sheets with a thickness of approximately 1.0 mm were prepared by the hot pressing method. The silicate layers of the clay were intercalated and randomly distributed in the matrix. The addition of OMMT to the PLLA/PCL blend significantly improved the tensile properties and dynamic mechanical properties of nanocomposites. On the other hand, layered silicate explicitly improve the thermal stability of PLLA/PCL blends when the OMMT content is less than 5 wt%. SEM images showed that adding OMMT could decrease the size of phase-separated particles, which made the material more uniform. More OMMT in nanocomposites brought about lower crystalline temperature and higher crystalline degree. Nanocomposites of isotactic polypropylene (iPP) and carbon nanoparticles (CN) were prepared using a twin-screw co-rotating extruder [65]. Nanocomposites containing 0.5, 1, 2.5 and 5 wt% CN were prepared by melt mixing in a twin-screw co-rotating extruder. Along the screw different screw elements existed which induced polymer melting and achieved fine dispersion of the nanoparticles in the polymer melt. The iPP pellets were fed into the throat of the extruder, while CN nanoparticles were introduced separately through a streamside feeding port further down the barrel directly into the polymer melt. Changing the feeding rate of each dosing unit automatically controlled the required proportions of both materials. The melt temperature was continuously recorded during compounding. After compounding, the material was extruded from a die, which had three cylindrical nozzles of 4 mm diameter, to produce cylindrical extrudates. Morphological characterization of the prepared materials was carried out by TEM and micro-Raman studies. At low loadings the filler was finely dispersed while at higher loadings there was a tendency for agglomeration.

5.5 Manufacturing Techniques of Polymer Gels A gel is a crosslinked polymer network swollen in a liquid medium. The properties of gels depend on the interaction of the polymer network and the liquid. The liquid prevents the polymer network from collapsing into a compact mass, and the network, in turn, retains the liquid. The phases that can be taken by a gel are sol, gas-like gel and liquid-like gel. At the critical point of the gel–gel phase transition, the two gel phases become indistinguishable. Polymerizing bifunctional and polyfunctional units through a condensation reaction can form a polymer network. The bifunctional units form linear chains and the polyfunctional units act as crosslinks. A polymer network can also be formed by crosslinking polymers formed from bifunctional monomers. A schematic representation of polymer gels is shown in Figure 5.9 [66]. Polymer gels can be classified as microgels, aerogels, xerogels and nanostructured gels, etc. 5.5.1

Microgels

Microgels are particles of crosslinked polymer networks dispersed in a medium, such as water. These gels can undergo rapid, reversible and dramatic phase transitions when triggered by slight environmental variations in temperature, pH or ionic strength. Medical applications for smart polymers range from diagnosis to intelligent

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Figure 5.9 A schematic of gel formation (studies on two types of built-in inhomogeneities for polymer gels: Frozen segmental concentration fluctuations and spatial distribution of cross-links. Reprinted from [66]. Copyright (2003) with permission from American Chemical Society.

distribution of drugs and actuators (such as artificial muscles and valves), while other sectors are benefiting from new technologies for microelectronic sensors and environmentally adaptive clothing. Figure 5.10 [67] shows the representation of agar microgels coated with a bi-layer of chitosan and fluprescein by Malverin U.K. Functionalized microgles were prepared and their properties have been investigated in detail [68]. In the preparation method, the functional monomer was dissolved in water and heated to the polymerization temperature of 70◦ C under a nitrogen purge. After 30 min, ammonium persulfate (APS) was dissolved in

Figure 5.10 Agar microgels coated with a bi-layer of chitosan and fluorescein labeled alginate imaged by confocal microscopy, Malvern, UK. (Posted on August 7th, 2009.) Reprinted from http://www.nanotech-now.com/ news.cgi?story_id=34180 Copyright (2011) Unilever, Colworth.

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water and injected to initiate the polymerization. Polymerization reaction was carried out in the prescence of ammonium persulfate (APS). Microgels were purified after cooling by ultracentrifugation, decantation, and redispersion in water and were lyophilized and stored at 4◦ C. Temperature-sensitive microgels were prepared by coplymerizing N-isopropylacrylamide (NIPAM) with triethyleneglycol methacrylate (TREGMA) containing long functional hydrophilic side chains (viz., -(-OCH2 CH2 -)3 -OH) [69]. The structure and morphology of the microgels obtained were investigated through 1 H NMR and transmission electron microscopy (TEM). Experimental results show that TREGMA can copolymerize well with NIPAM to prepare microgels with regular spherical morphology. Investigation on the volume phase transition of the microgels through turbidimetric method, dynamic light scattering (DLS) and differential scanning calorimetry (DSC) revealed that the incorporation of hydrophilic side chains -(-OCH2 CH2 -)3 -OH has strong influence on the sizes, size distribution and the deswelling properties of the microgels. The microgels exhibit much larger deswelling ratios (α) than pure poly(N-isopropylacrylamide) (pNIPAM) microgels.

5.5.2

Aerogels

Aerogels are dried gels with a very high relative pore volume. These are versatile materials that are synthesized in a first step by low-temperature traditional sol-gel chemistry. As aerogels combine the properties of being highly divided solids with their metastable character, they can develop very attractive physical and chemical properties not achievable by other means of low temperature soft chemical synthesis. Aerogels of PVDF-HFP were prepared by dissolving the polymer in acetone followed by ethanol at different concentrations [70]. Each polymer solution was placed in a formation cell. The formation cell was reposed in a freezer at –20◦ C until a gel was formed. The gel was placed in the vessel, which was closed and filled from the bottom with super-critical carbon dioxide (SC-CO2 ) up to the desired pressure using a highpressure pump. The membrane-like aerogels exhibited modulated pore size and porosity. These morphological achievements are responsible for the enhanced transport and surface properties of the membranes. In addition, these membrane-like aerogels exhibit polymorphic character, the crystalline composition of which could be usefully changed as a function of the solvent/nonsolvent ratio used. Incorporation of 3-aminopropyltriethoxysilane into the silicate sol-gel process provides functionality on the backbone of the silica aerogel and, that promotes grafting of various polymers (polyisocyanates, epoxides, etc.) into the aerogel structure [71]. The intimate mixing of the polymer crosslinks within the aerogel structure and the covalent bonding between the inorganic and organic phases are essential in controlling the material properties.

5.5.3

Xerogels

A xerogel is a solid formed from a gel by drying with unhindered shrinkage. Xerogels usually retain high porosity, and enormous surface area, along with very small pore size. When solvent removal occurs under hypercritical conditions, the network does not shrink and a highly porous, low-density material known as an aerogel is produced. Heat treatment of a xerogel at elevated temperature produces viscous sintering and effectively transforms the porous gel into a dense glass. The sol-gel process offers a convenient and efficient method for the polymerization of alkoxysilanes xerogels under extremely mild conditions. One of the methods for preparing xerogels based on alkoxysilane is described as follows [72]. Solvent and catalyst are added to the monomer solution in a volumetric flask. The solution is shaken vigorously to ensure homogeneity, poured into a polyethylene bottle. Gellation normally occurs within 1–12 hours. The gel is then aged for one week and processed. The materials are

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processed by either solvent series or water processing conditions, air-dried for 2–3 days, ground, and dried under dynamic vacuum at 100◦ C to afford the corresponding amorphous polysilsesquioxane xerogel in approximately quantitative yield. Bridged polysilsesquioxanes have been applied for a number of applications including coupling agents for composites, high-surface-area catalytic materials, and protective coatings. These materials have also shown considerable promise as chiral catalysts, metal scavengers, chromatographic solid phases, membrane materials, and photonic and electronic materials. 5.5.4

Nanostructured Gels

Nanostructured organic polymer gels are promising materials for a wide range of applications such as supports for catalysts, sorbents, reactive membranes, chromatography, drug delivery, or as components of nanocomposite materials. Compared to their inorganic counterparts they typically exhibit different mechanical and chemical properties such as elasticity or mechanical and chemical resistance, and they can be easily functionalized by copolymerization or polymer-analogous reactions. A few examples of nanostrucrured gels are cited below. Free radical polymerization of divinyl benzene (DVB) and styrene (St) was performed within reverse hexagonal phases of the anionic surfactant bis(2ethylhexyl) sulfosuccinate sodium salt (AOT) [73]. The morphologies before and after polymerization were studied by small angle X-ray scattering (SAXS), polarized light microscopy, and scanning and transmission electron microscopy. Polarized light microscopy and SAXS showed no significant changes in textures and order during polymerization. A detailed investigation of the polymer matrix after polymerization and removal of the template revealed the formation of regular micrometersized layerlike morphologies. The polymer layers had a hierarchical structure, as they consisted of aligned strings of polymer beads with diameters of about 100 nm. Structure formation was directed by the lyotropic phase during polymerization-induced phase separation on a nanometer scale. Nanocomposite gels (NC gels), consisting of poly- (N-isopropylacrylamide) (PNIPA) and inorganic clay (hectorite), prepared using a wide range of clay concentration has been reported [74]. Fine, uniform dispersions of all constituents were achieved by adopting an improved mixing procedure utilizing rotation and revolution. The rotation/revolution mixing was carried out effectively in two stages such as (1) clay/NIPA aqueous solutions and (2) mixing KPS and TEMED into the solutions. The latter was conducted at low temperature to suppress the initiation reaction, before moving the reaction into a water bath at 20◦ C to initiate the in situ free-radical polymerizations. Uniform hydrogels were obtained in both cases (NC gels and OR gels) after maintaining the reaction at 20◦ C for 20 h. All NC gels were uniform and transparent, almost independent of the clay content. The effects of clay content on the tensile mechanical properties on the first and second cycles, the time-dependent recovery from the first large elongation and the optical anisotropy of NC gels, the disappearance of the glass transition and the formation of clay-polymer intercalation in the dried NC gel were revealed. Thus, it became clear that the properties and the structure changed dramatically for an NC gel with critical clay content. 5.5.5

Topological Networks

A novel class of gel materials known as a topological network, characterized by the sliding character of the crosslink points (also called slide-ring crosslink points), has been recently proposed. The topological network architecture is obtained by the intermolecular crosslinking reaction of precursor polyrotaxanes. The main characteristic of the sliding crosslink points is to allow a sliding motion of the constitutive template network chains through the figure-of-eight junctions. The sliding network architecture was experimentally realized by the use of high molecular weight polyrotaxanes, with chemical structures based on R-cyclodextrin macrocycles (R-CDs) threaded along a template poly(ethylene glycol) (PEG) and their intermolecular bridging

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Figure 5.11 Schematic representation of the molecular architecture of a sliding gel obtained by reaction between R-CDs/PEG polyrotaxanes and divinyl sulfone. Reprinted from [75]. Copyright (2007) with permission from American Chemical Society.

with a crosslinker through the hydroxyl groups of the R-CDs, forming the crosslinking network points in the form of a figure-of-eight. The sliding gels are expected to have unusual chemical, physical and mechanical properties because of the theoretical ability of the crosslinking points to slide along the template polymer chain. Topological polymer networks with sliding crosslink points known as ‘sliding gels are prepared’ as a new class of supramolecular networks based on intermolecularly crosslinked R-cyclodextrins/ poly(ethylene glycol) polyrotaxane precursors. The scheme of preparation is shown in Figure 5.11[75]. The crosslink points of such networks are not fixed but can slide along the template chain of the polyrotaxanes. The main parameters governing the sliding gel properties are the number of cyclodextrins per polyrotaxane, the crosslinking density and the nature of the swelling solvent. Small angle neutron scattering, swelling measurements and mechanical spectrometry are used to understand the unusual physical properties and their relation to the molecular structure of the sliding gels. The swelling as well as the viscoelastic properties is found to be solvent dependent reflecting the structural changes of the network. Indeed, in water, the number of crosslink points (topological and physical) increase as opposed to dimethyl sulfoxide.

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Hydrogels

Hydrogel (aquagel) is a network of polymer chains that are water-insoluble, sometimes found as a colloidal gel in which water is the dispersion medium. Hydrogels are highly absorbent natural or synthetic polymers. Hydrogels also possess a degree of flexibility very similar to natural tissue, due to their significant water content. A few methods of prepartion of hydrogels are narrated here. Methacrylate-structured poly (ampholyte)s were synthesized in the homopolymer and copolymer forms starting from the N-methacryloyl-L-histidine (MHist) and the N-isopropylacrylamide (NIPAAm) and also obtained in the hydrogel form. They exhibited close thermodynamic behavior and the viscometric data revealed that the minimum hydrodynamic volume of the polymer at its isoelectric point (pH 5) shifted to lower pHs as the NIPAAm content increased, and beyond a critical low MHist content the reduced viscosity decreased, even at low pHs [76]. The phenomenon was attributed to hydrophobic forces between the isopropyl groups outweighing the repulsive electrostatic interactions of the polymer in the positively charged form. The effects of crosslinker contents on various physical properties of Poly (N-isopropylacrylamide) (PNIPA) hydrogels, i.e., nanocomposite type PNIPA hydrogels (NC gel) and conventional chemically crosslinked PNIPA hydrogels (OR gel) were investigated [77]. While NC gels were composed of a unique organic PNIPA)/inorganic (clay) network, the inorganic clay acted as a multifunctional crosslinker in place of an organic crosslinker (BIS) as used in OR gels. NC and OR gels were prepared using initial solutions consisting of monomer (NIPA)], crosslinker (clay or N, N-methylenebis (acrylamide) (BIS), solvent (H2 O), initiator (potassium peroxodisulfate (KPS) and catalyst (tetramethylenediamine (TEMED). First, a transparent aqueous solution consisting of water, inorganic clay and NIPA was prepared. Next, the catalyst and subsequently the aqueous solution of initiator were added into the former solution with stirring at iced-water temperature. Then, free-radical polymerization was allowed to proceed in a water bath. In the preparation of OR gels, BIS was used as a crosslinker. Throughout the experiments, oxygen was excluded from the system. Both NC and OR gels were synthesized in two kinds of glass vessels, one a columnar shape and the other a rod shape. NC gels exhibited extraordinary mechanical toughness, the tensile moduli and tensile strengths are almost proportional to the clay content, while the elongation at break tended to decrease slightly with increasing clay. On other hand, OR gels were weak and brittle natured and there was no detectable change in properties on altering the concentration of BIS (CBIS). PNIPA/clay nanocomposite gels with co-crosslinked networks were synthesized by in situ, free radical polymerization of NIPA (N-isopropylacrylmaide) in the presence of two types of crosslinker, an inorganic crosslinker (clay: hectorite) and an organic crosslinker (N, N-methylenebis(acrylamide)(BIS), with concentrations n and m, respectively, in aqueous media [78]. The optical properties and the tensile and compressive mechanical properties of the resulting hydrogels (NCn-ORm gels) were investigated and are discussed herein in terms of co-crosslinked PNIPA network structures. NCn-ORm gels were all uniform, but their transparencies changed considerably according to n and m and were generally different from the sum of the transparencies of corresponding NCn and ORm gels. NCn-ORm gels generally exhibited pronounced weakness and brittleness in tensile tests, like ORm gels. In contrast, in compressive mechanical tests, large improvements were achieved at high n and low m values. Furthermore, abnormal increases in modulus were observed in both mechanical tests. The properties of the gels were mainly due to the formation of a ‘microcomplex structure’ consisting of exfoliated clay platelets and PNIPA chains with enhanced chemical crosslinking. PNIPA-NC gels (abbreviated to N-NC gels) were prepared by in situ, free-radical polymerization of NIPA monomer in the presence of exfoliated clay uniformly dispersed in aqueous media. N-NC gel films of 2 mm thickness were prepared in a laboratory-made molding apparatus made from poly(methyl methacrylate) substrates [79]. Conventional PNIPA-OR gels (N-OR gels) and viscous solutions of linear PNIPA (N-LR) were also prepared using the same procedure, with and without organic crosslinker (BIS), instead of clay. Extraordinarily high hydrophobicity was observed on the surfaces of PNIPA nanocomposite hydrogels (N-NC gels). The high hydrophobicity of N-NC gels may be caused by the effective alignment of N-isopropyl groups

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of PNIPA chain at the gel-air interface and is enhanced by the other factors, such as network structure, water content, and topological roughness. The surface of N-NC gels exhibited hydrophobic-to-hydrophilic changes by changing the circumstance from dry (in air) to wet (in water), and vice versa. Poly(lactide)-block-poly (ethylene oxide)-block-poly(lactide) (PLA-PEO-PLA) triblock copolymers are known to form physical hydrogels in water as a result of the polymer’s amphiphilicity [80]. Their mechanical properties, biocompatibility, and biodegradability have made them attractive for use as soft tissue scaffolds. However, the network junction points are not covalently crosslinked, and in a highly aqueous environment of these hydrogels adsorb more water, transform from gel to sol and lose the designed mechanical properties. Hydrogels with varying crosslinking ratio and ionic content were prepared from interpenetrating networks of poly (vinyl alcohol) and poly(acrylic acid [81]). Hydrogels with large Mc values were found to swell to a greater extent than those with small Mc values. It was also observed that an increase in Mc values yielded faster swelling and deswelling rates, as the rates for membranes with Mc = 18,000 were about twice as fast as were the rates for membranes with Mc = 34,000. Oscillatory swelling behavior was investigated in response to changes in the pH and ionic strength of the swelling medium. A change in pH from 3 to 6 caused an ionization of the hydrogels and an increase in the weight swelling ratio, with a greater increase exhibited by IPNs with a higher ionic content. Increase in pH also caused an increase in the average mesh size. 5.5.6.1

Double Network Hydrogels

Double network gels (DN gels) are recently-developed high strength gels in which a second network is polymerized without any crosslinkers [82]. Truly Independent-DN gels (named ‘t-DN’ gels), which do not have any covalent bonds between the first and the second networks, and normal double network gels, named as c-gels, have the high strength of usual DN gels without the crosslinker of the second network is actually achieved by the interconnection between the two networks through covalent bonds. Further, it was found that the t-DN gels became stronger than the c-DN gels when the second network was loosely crosslinked. Both t- and c-DN gels (normal double network gels) were prepared by two-step free-radical polymerization. The preparation routes and chemical structures of both the DN gels are shown in Figure 5.12 [82]. The first network active-PAMPS (poly (2-acrylamido-2-methylpropanesulfonic acid) gels (named ‘a-PAMPS’gels) were synthesized using AMPS as monomer, N, N - methylenebis(acrylamide) (MBAA) as crosslinker, and 2-oxoglutaric acid as UV initiator. To perform the polymerization, the solution was purged in an argon atmosphere to remove dissolved oxygen, and generated and almost all of the unreacted double bonds remaining in the PAMPS gels reacted to form an inert group. Thus, inert-PAMPS gels were prepared (named ‘i-PAMPS’ gels). The c- and t-DN gels were synthesized from the a- and i-PAMPS gels, respectively. Synthesized DN gels were immersed in a large amount of pure water at least for 5 days to remove residual unreacted reagents. All photopolymerizations were performed under argon. c-DN gels are connected to the first network by copolymerizing with the unreacted crosslinker of the first network, which forms internetwork crosslinking (INC) points. This structure explains the anomalously high strength of the c-DN gels without adding any crosslinkers for the second network. That is, the force applied to the DN gel is transferred from the first network to the second one through this interconnected structure. On the other hand, the t-DN gels, which consist of two truly independent and interpenetrated network structures, also show a high toughness when the crosslinker density of the second network is just above the lowest critical and optimal value

5.6 Manufacturing Techniques of Interpenetrating Polymer Networks (IPNs) An interpenetrating polymer network (IPN) is an intimate combination of two polymers both in network form, at least one of which is polymerized or crosslinked in the immediate presence of the other. An IPN is

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Figure 5.12 Chemical structures and the mechanical properties of active-PAMPS gels, inert-PAMPS gels, connected-DN gels, and truly independent-DN gels. Reprinted from [82]. Copyright (2009) with permission from American Chemical Society.

distinguished from simple polymer blends, blocks, or grafts in two ways: it swells but it does not dissolve in solvents, and the creep and flow are suppressed. There are different types of IPNs, such as full IPN, sequential IPN (SIPN), simultaneous IPN (SIN), thermoplastic IPN, semi IIPN, semi-II IPN, and pseudo IPN. A schematic representation of an ideal IPN and a hybrid IPN is shown in Figure 5.13 [83].

5.6.1

Full IPNs

A full IPN comprises of two networks that are ideally juxaposed, which generate a lot of entanglements and interaction between the networks [84]. A full IPN of crosslinked gelatin and polyacrylonitrile was prepared using glutaraldehyde and N, N’methylene bisacrylamide as the crosslinking agents respectively [85]. This resulted in getting a hard biomaterial that displayed water uptake potential. It was found that with increasing concentrations of gelatin, PAN and MBA, the IPN showed a decreasing tendency of water sorption. The IPNs exhibited a fair degree of microhardness, which increased with increasing the content of gelatin, PAN and MBA in the IPN. This was also confirmed with the thermal studies, which revealed the increase in the value of glass transition temperature of IPNs. The in-vitro blood compatibility of the prepared IPNs also depended on the chemical architecture of the network. The prepared IPNs were also assessed for in-vitro blood compatibility by methods such as protein (BSA) adsorption, blood-clot formation and percent hemolysis measurements. Homo-IPNs [84] are a special type of full IPNs, where both polymers used in the networks are the same. They are usually sequential IPNs. They are one of the first types of IPNs to be commercialized and they are used as model materials for theoretical work.

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(b)

Figure 5.13 (a) Ideal IPN structure where the two networks are fully developed and interpenetrated. (b) More realistic structure of a hybrid IPN, in which the hard inorganic network domains are not necessarily connected but rather dispersed in the soft organic polymer matrix. The organic network is represented by the hard black lines. The ZrO2 network is represented by the lighter lines. Reprinted from [83]. Copyright (2006) with permission from Macmillan Publishers Ltd.

5.6.2

Sequential IPNs

An IPN is formed by polymerizing the first mixture of monomer, crosslinking agent, and initiator or catalyst to form a network. The network is swollen with the second combination of monomer and crosslinking agent and polymerized to form an IPN [84]. Natural rubber/PS full and semi-IPNs were prepared by sequential technique [86]. The NR was cross linked using DCP followed by polymerization and crosslinking of PS phase. In a typical method the vulcanized NR sheets were weighed and kept immersed in an inhibitor-free monomer containing benzoyl perodixe as initiator and divinyl benzene as the crosslinker for the PS phase. The NR samples were swollen to different time intervals to obtain varying weight percentages of PS. The swollen samples were kept at 0◦ C for a few hours to achieve equilibrium distribution of styrene monomer in the matrix. The swollen networks were heated at 80◦ C for 6 h and at 100◦ C for 2 h to complete the polymerization and crosslinking of styrene monomer. The hardened sheets were kept in a vacuum air oven to make them free of unreacted styrene, the final weight of the IPN was taken and composition of the sample determined (Figure 5.14) [86]. The mechanical performance IPNs improved upon increasing the PS crosslinking. Improvement was observed in tensile strength, tear strength and hardness on increasing the crosslinker content of PS up to 4. On increasing the PS crosslinking up to 4% DVB, the phase distribution became more and more uniform. Interpenetrating networks of poly (acrylic acid) (PAAc) and poly (N-isopropylacrylamide) (PNIPAAm) were synthesized in two consecutive steps utilizing ionizing radiation in the first step and chemical reaction in the second step [87]. The first network of PAAc hydrogel was formed by ionizing radiation (gamma or electron beam). The secondary network of PNIPAAm was synthesized directly within the primary network of PAAc from an aqueous solution of N-isopropylacrylamide (NIPAAm) containing a crosslinking agent, accelerator and redox initiator. The interpenetrating networks (IPNs) were characterized morphologically by scanning electron microscopy (SEM), and their thermal behavior was analyzed by differential scanning calorimetry (DSC) and thermogravimetry analysis (TGA).

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polymerisation (0% DVB)

D0, A0, B0 series

swollen NR network

NR network

polymerisation crosslinking (2% DVB)

D1, A1, B1 series

polymerisation crosslinking (4% DVB)

NR network

swollen NR network

D2, A2, B2 series

styrene monemer Inlletor crosslinker NR Chain

polymerisation crosslinking (6% DVB)

PS Chain D3, A3, B3 series

Figure 5.14

5.6.3

A schematic representation of Formation of NR/PS IPN. Reprinted from [86].

Simultaneous Interpenetrating Networks (SINs)

An IPN is formed by polymerizing two different monomer and crosslinking agent pairs together in one step. In this process, the two components are polymerized by reactions that will not interfere with one another. This occurs by polymerizing one network by a condensation reaction, while the other network is formed by a free radical reaction [84]. IPNs formed by a simultaneous reaction of components are inherently simpler to make since there is only one reaction required to make the IPN while the sequential method requires two reactions. If the monomers are chosen correctly, SINs can be injection molded, which is a real advantage over the sequential IPNs. Sequential IPNs must be either compression molded or synthesized in bead form. The swelling of the network in the

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sequential method can get a little messy too. These factors make simultaneous IPN reactions a little more appealing to people who are looking to commercialize IPN materials [84]. A series of castor oil polyurethane/poly (methyl methacrylate) interpenetrating polymer networks (IPNs) curing at room temperature were synthesized by a simultaneous method [88]. Component analysis and viewing results of morphology and miscibility among the multi-phases of materials obtained by TEM indicated that the systems belong to graft-mode IPNs and domains between two phases were controlled in the range of nanometer scale. New biodegradable hydrophobic polyurethane (PU)/hydrophilic poly (ethylene glycol) diacrylate (PEGDA) IPN was simultaneously synthesized with changing the molecular weight of PEGDA to investigate the effect of crosslinking density on the degree of phase separation [89]. PU was modified using biodegradable poly (ε-caprolactone)diol and the hydroxy group of PEG was substituted to crosslinkable acrylate group having double bond, which induce photo-polymerization. Crosslinkable PEGDA was prepared by substituting the hydroxyl end group of PEG by acrylate group. Acrylolyl chloride dissolved in benzene, in three fold molar excess based on hydroxyl groups of PEG, was dropped to the PEG solution having TEA at room temperature. The reaction proceeded overnight at 35◦ C under nitrogen. After removing triethylamine hydrochloride salt, the filtrate was precipitated in hexane twice. The precipitated PEGDA was dried at 40◦ C for 24h under vacuum. 5.6.4

Latex IPNs

Latex IPNs are formed by an emulsion polymerization technique. The morphology of the IPN depends upon how the IPN components are polymerized. Both monomers can be added at once, which will tend to give a more uniform morphology in the particles (a simultaneous IPN formation). The monomers can also be added in stages. For instance, monomer 1 can be polymerized to form latex and monomer 2 can then be added (a sequential IPN formation). Depending on how fast monomer 2 diffuses into the latex, one can get either a homogeneous incorporation of the monomer in the latex or most of monomer 2 may react near the surface of the latex particle. If most of monomer 2 reacts near the surface, one has a core-shell morphology. The finished latex particles can be collected and used as they are or they can be used as a coating. For instance, a series of latex particles with interpenetrating polymer network structure have been synthesized from waterborne polyurethane (PU) and polystyrene (PS) [90]. The effect of PU/PS composition, crosslinking density in the PS domain as well as in PU have been studied in terms of dispersion size, transmission electron microscopy morphology, mechanical and dynamic mechanical properties, in addition to swellability in water and toluene of the dispersion cast film. It was found that inverted core (PS)–shell (PU) morphology was well defined and that both the domain size and the film properties were well controlled by the latex composition and crosslinking density of both phases Functional poly[styrene-co-(2-ethylhexyl acrylate)-co-(1,6-hexanediol diacrylate)]/poly[(methyl methacrylate)-co-(butyl acrylate)-co-(methacrylic acid)-co-(diacetone acrylamide)] (PS/PA) [91] semiinterpenetrating polymer networks (semi-LIPNs) containing ketone carboxyl groups were synthesized by two-stage emulsion polymerization, and characterized by Fourier transform infrared, transmission electron microscopy, dynamic light scattering and DSC. A unique feature of the PS/PA semi-LIPNs is their ability to crosslink and form thermosetting full-IPN polymers through the reaction of ketone carboxyl and hydrazide in the course of film formation at ambient temperatures. Series of latex particles with various levels of crosslinking density in the PS and PA domain and varied composition of PS/PA LIPNs were obtained. The effects of the LIPN PS/PA composition and the level of crosslinking density in the PS and PA domain on film density, swell ratio, mechanical properties and contact angle with water were investigated. Maximum synergy effects obtained at around 50/50 (PS/PA) in terms of mechanical properties, density and contact angle with water, indicated that the maximum degree of interpenetration is obtained at this composition. A series of

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poly(styrene-acrylonitrile)/poly(ethyl acrylate-n-butyl acrylate) latex interpenetrating polymer networks (LIPNs) were synthesized by changing the kind of crosslinker and introducing a buffer [92]. The results show that the crosslinker has an important effect on the damping properties of the LIPNs; divinylbenzene is the best crosslinker in the study. Moreover, introducing a buffer into LIPNs leads to an increase of the damping values over the temperature range where the damping value surpasses 0.5. The LIPN blend prepared by mixing LIPNs results in broadening of the damping peak, therefore improving the damping properties. 5.6.5

Thermoplastic IPNs

Thermoplastic IPNs consists of a block copolymer like SBS rubber and a semi-crystalline or glassy polymer. Thermoplastic IPNs are moldable, and can be extruded, and recycled. These IPNs have physical crosslinks rather than chemical crosslinks like, thermoplastic elastomers. Typical physical crosslinks arise from ionic groups, crystallinity, or glassy domains. The structure-property relationships of thermoplastic polymer blends based on poly(ether-urethane) ionomer (PEUI) and K+-Ion containing styrene-acrylic acid random copolymer (S-co-AA(K)) have been investigated by different methods [93]. Convergence of the glass transition temperature values of the PEUI and the S-co-AA(K) components in the blends was studied and improvements of end-use properties have been found. Thermoplastic interpenetrating polymer networks (t-IPNs), prepared by melting and pressing of crystallizable polyurethane (CPU) and styrene/acrylic acid random copolymer (S/AA) in wide ranges of composition, were investigated by the combination of various thermal analysis techniques: differential scanning calorimetry (DSC); thermomechanical analysis (TMA); thermally stimulated depolarization currents (TSDC); and thermally stimulated conductivity (TSC) measurements, as well as broadband dielectric relaxation were microheterogeneous systems with contributions to microheterogeneity from both the heterogeneity of the individual polymers and the thermodynamic incompatibility of the components. 5.6.6

Semi-IPNs

Interpenetrating polymer networks of poly(2-hydroxyethyl methacrylate) and poly(ethylene oxide) was synthesized by the free radical copolymerization of 2-hydroxyethyl methacrylate and ethylene glycol dimethacrylate in aqueous solutions of poly(ethylene oxide) diglycidyl ether (PEODGE), followed by amine−epoxy coupling reactions within the resultant gel network [94]. The effects of ethylenediamine and PEODGE concentrations and coupling methods were examined. Swelling studies and cryogenic-SEM analysis showed that these novel networks have high water intakes and pore sizes on the order of up to 1 μm. A semi-interpenetrating polymer network (IPN) formed from polyurethane, starch and polymethyl methacrylate (PMMA) was prepared by varying the crosslinking agent, initiator, and molar ratio of the monomers [95]. IPNs began to synthesize, substituting the polyol OHs for functional groups of starch. The starch concentration was 3 wt.%. The synthesized semi-IPNs were evaluated by means of dynamic mechanical analysis (DMA). It was found that a relationship existed between the quantity of crosslinker and the storage modulus, where alpha and beta relaxations were present in the polymethyl methacrylate’s rigid chains, and this peak widened upon increasing the quantity of PMMA in the semi-IPN. The morphology of the semi-IPNs shown by SEM and AFM micrographs suggests that 0.5 wt.% of initiator and 3 wt.% of crosslinker demonstrate uniform surfaces. Three-dimensional micrograph of semi-IPNs (Figure 5.15) by atomic force microscopy shows a regular lamella formation between 232 and 63.35 nm in average height. This lamella was attributable to a polyurethane IPN, and the 250 nm average height lamellas with a white form were attributable to poly(methyl methacrylate). These results were in good agreement with SEM results and demonstrated that a structure like an interweave was obtained.

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µm/div 0.13

0

2.0 µm/div

2.0 µm/div

Figure 5.15 AFM micrographs of a 50=50 PU=PMMA semi-IPN using 0.5 wt.% initiator (BPO) and 3% crosslinker (TRIM). Reprinted from [96]. Copyright (2009) with permission from Taylor and Francis.

Semi-IPNs based on PU were produced by sequential method [96]. In the first stage, synthesis of crosslinked polyurethane (PU) from poly(oxypropylene glycol) and triisocyanate adduct (2,4-, 2,6 toluene diisocyanate and trimethylol propane) was carried out at 60◦ C up to full conversion of the functional groups (monitored by IR-spectroscopy). Synthesized products were dried under vacuum up to constant weight. In the second stage, the second component was introduced into the network by swelling it in the monomer (butylmethacrylate (BMA) and styrene (St)). Initiator of radical polymerization of monomers (2,20-azo-bis-isobutyronitrile) was dissolved in monomers before swelling. The amount of monomer introduced into the network by swelling was dependent on the crosslinking density; the identical conditions of reaction in each case were supported. The amount of the second component in the network depended on the network density. In such a way semi-IPNs obtained using networks with different crosslinking density had various ratios of components. The amount of a monomer in the network was determined by the gravimetric method. Alginate and alginate:chitosan semi interpenetrating polymer network (IPN) scaffolds were prepared by the freeze-drying process [97]. Alginate scaffolds were crosslinked with different concentrations of CaCl2 , after freeze-drying. Scanning electron microscope (SEM)/Energy Dispersive Analysis by X-ray (EDAX) analysis and swelling studies indicated that crosslinking of scaffolds with 3% CaCl2 had effectively created suitable alginate scaffolds in terms of optimum porosity and mechanical stability. The structural and cellular outcomes demonstrate potential utility of chitosan semi IPNs in alginate scaffolds. Semi-interpenetrating polymer networks (semi-IPNs) based on polyvinyl alcohol (PVA) and crosslinked polyacrylamide (PAM) were prepared by redox polymerization [98]. Various specimens were prepared by varying concentration of PVA, acrylamide (AM) and crosslinker MBA to study the effect of different compositions on the structural and mechanical property of semi-IPNs. The structural and morphological characterizations of prepared semi-IPNs were obtained from the studies of FTIR spectroscopy, X-ray diffraction (XRD) and environmental scanning electron microscopy (ESEM). The mechanical properties of pure PVA and semi-IPNs like tensile strength, percentage elongation and deformation under stress were obtained by load-displacement curves. A similar profile for deformation was obtained for all the semi-IPNs. It was found that the elastic modulus, necking behavior and ultimate failure were largely affected by the chemical composition of the semi-IPNs.

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Pseudo-IPNs

Synthesis of amphiphilic pseudo-semi-interpenetrating polymer networks (pseudo-semi-IPNs) containing linear poly(styrene) and poly(ethylene glycol) (PEG) has been crosslinked through monodendritic fragments [99]. A unique feature of the synthetic strategy is the permanent attachment of the linear segment to the PEG network by a transesterification reaction between the hydroxyl groups at both ends of the PEG and peripheral ethyl ester moieties in the monodendron portion of a linear poly(styrene)−dendritic poly(benzyl ether) AB block copolymer. The proceeding of the reaction is monitored by 1 H NMR and size exclusion chromatography (SEC). The formation of an interlock structure between the linear block and the network matrix in the pseudo-semi-IPN is evidenced by the results from spectroscopic analyses and differential scanning calorimetry measurements. The accessibility of functional centers in the grafted semi-IPN is confirmed by model reactions with fluorescent markers, fluorescence spectroscopy, and NMR techniques and shows the potential of these novel materials as sequestering reagents for resin capture−release applications in parallel synthesis, combinatorial chemistry and advanced drug design. Pseudo-interpenetrating polymer networks (p-IPNs) of an epoxy resin and poly(methyl methacrylate) (PMMA) were synthesized simultaneously [100]. Methyl methacrylate (MMA) was polymerized with benzoyl peroxide in the presence of an oligomeric epoxy resin, which contained stoichiometric amounts of either a tetrafunctional (i.e., ethylenediamine (EDA)) or hexafunctional (i.e. triethylenetetramine (TETA)) crosslinking agent. The resulting materials would be useful for aerospace, automotive, medical, and dental applications. The structure of the epoxy resin (i.e., tetraglycidyl 4,4-diamine diphenyl methane) was confirmed by elemental analysis and IR and NMR spectroscopies. The thermal, mechanical, and kinetic behaviors of the polymeric networks produced from this oligomeric epoxy resin and the p-IPNs of this crosslinked epoxy with linear PMMA were compared [101]. Time-resolved IR spectroscopy of the cyclic ether group in the epoxy at 906 cm−1 revealed kinetic information between 40 and 80◦ C for a complex reaction scheme. The thermal expansion coefficients of the crosslinked epoxy resin increased at higher concentrations of PMMA, but hexafunctional TETA produced a higher crosslink density and reduced the thermal expansion relative to the same networks and p-IPNs prepared with tetrafunctional EDA. Elastic modulus, fracture stress and toughness increased synergistically when the epoxy resin was crosslinked with EDA in the presence of 10 wt % MMA. This p-IPN exhibited minimal shrinkage, on the order of 0.15%, during the actual polymerization. PDMS-PMMA-based interpenetrating polymer networks gained attention recently [102]. Firstly, the PDMS elastomer was prepared in the presence of MPS, and the polymerization reaction was catalyzed by H2 SO4 . The experimental results indicated that almost all the MPS had been hydrolyzed and incorporated with PDMS. Secondly, the MMA monomer was introduced in the networks of PDMS, through a further polymerization and the PDMS-PMMA IPNs, without obvious phase separation, could finally be produced.

5.7

Conclusion and Future Outlook

An overview of various manufacturing techniques of multiphase polymeric systems is discussed. Polymer blends are mainly prepared by solution and melt blending techniques. Among the different types of blending techniques, melt blending is the most common method for polymer blends. The control of phase morphology is more authentic in the melt blending process. The shape, size and spatial distribution of the phases result from a complex interplay between viscosity (and elasticity) of the phases, interfacial properties, blend composition and processing parameters. The real time (on-line) monitoring of morphology of polymer blends during the manufacture and morphology control are current fields, which require a lot of attention. Composites are prepared by various techniques according to their end use. Even though the hand layup process is operator-dependent and labor intensive, it is attractive due to its low capital outlay, flexibility and

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lack of size limit. Along with the traditional single-vacuum-bag process, a double-vacuum-bag process is also in use. Volatile management is more effective in the DVB process. Resin transfer molding is a commonly-used method for the manufacturing of fiber-reinforced thermosetting composites materials. RTM has its ability to encapsulate metal ribs, stiffeners, inserts and so on within the molded laminate. The filament winding is a noteworthy method for the processing of reinforced plastic storage tanks and pipes for chemical and other industries. The main advantages are: excellent mechanical properties; high production rate; good control of fiber orientation; good thickness control; good fiber content control; and good internal finish. Reaction injection molding (RIM) has emerged as a popular method of producing premium quality molded polymeric products. Different types of composite systems such as foam or solid, rigid or elastomeric, can be molded with fiber reinforcements through the RIM process. Overall improvement in stiffness and impact strength are enhanced by adding reinforcement in the material stream, or by using a molded preform in the mold that is encapsulated. Polymer nanocomposites with different nanofillers (nanoclays, nanotubes, nanowires and nanoparticles) have been developed recently with conventional methods such as solution intercalation, in situ polymerization and melt intercalation methods. The dispersion of nanoparticles is still a major problem facing nanocomposite research. In situ monitoring of flow and structure of nanocomposites during fabrication is an important area where a lot of research has to be undertaken. Synthesis of different kinds of polymer gels such as microgels, areogels, hydrogels, and nanostructured gels have been discussed. All these gels have potential applications in several fields such as artificial organs for medical treatments, soft contact lens, superabsorbent polymeric gels used in sanitary napkins and disposable diapers, carriers for protein and nucleic acid in gel electrophoresis, amendments in greening and agriculture, refrigerants in medical and food industries, etc. Hydrogels are prepared based on several polymers but poly-(N-isopropylacrylamide) (PNIPA) is more preferred because of its stimuli sensitivity. Incorporating PNIPA with nanoparticles makes it an excellent functional material. A new class of polymeric gels, known as topological gels, has been developed. The unusual properties of this new class of topological materials are decided by the complexation degree, the crosslinker fraction, and the interactions between the swelling solvent and the constitutive parts of the network. Finally the manufacturing techniques of different types of IPNS such as full IPNs, sequential IPNS, latex IPNs, thermoplastic IPNs, semi IPNs and pseudo IPNs have been discussed.

References 1. P. K. Malllick, Manufacturing. Chapter 5 in Fiber Reinforced Composites, Materials, Manufacturing and Design, Marcell Dekker, INC 270 Madison Avenue, New York (1988). 2. P. Ball, Techniques of manufacture of composites, Chapter 2 in Handbook of Polymer Composites for Engineers, Leonard Hollaway (Ed.), Jaico Publishing House, Bombay (1995). 3. D. H. M. Jones, Polymer Processing, Chapman and Hall Ltd, London, (1989). 4. G. Lubin (Ed.), Handbook of Composites, Van Nostrand Reinhold, New York, (1982). 5. L. A Utracki, ‘Mixing, blending and compounding’, Int. Plastics Engineering and Technology, 3, 1–46 (1999). 6. R. Asaletha, M.G. Kumaran, S. Thomas, Euro. Polym. J, 35, 253–271 (1999). 7. Li Zhenga, Yu-Z Wanga, Ke-Ke Yanga, Xiu-Li Wanga, Si-C Chena and J Lib, Euro. Polym. J, 41 1243–1250 (2005). 8. E. L. Bedia, S. Murakami , K. Senoo, S. Kohjiya, Polymer, 43, 749–755 (2002). 9. W-h Sun, X-p Qiao, Q-k Cao and J-p Liu, Chin. J. Chem. Phys., 23, 107 (2010). 10. T. M. M. Swamy, Siddaramaiah, J. H. Lee, Miscibility behavior of polyethylene oxide/hydroxypropyl methyl cellulose blends Advanced Materials Research, 47–50, 234–237 (2008). 11. S. S. Ochigbo, Æ Adrian, S. Luyt, Æ W. W. Fockeref, J Mater Sci, 44: 3248–3254, (2009). 12. N. Moonprasith, K. Suchiva, and O. Tongcher, M.Tech Thesis, National Metal and Materials Technology Center, Thailand.

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13. C. C. Ho, M. C. Khew, Y. F. Liew, Surface and Interface Analysis, 32, 133–143 (2001). 14. M. T. Shaw, Ch. 4 in Polymer Blends and Mixtures, series E, No: 89, D. J. Walsh, J. S. Higgins and A. Macennactoic (Eds), Marltime Nijhoft (pub) Dodrecht (1985). 15. C-T Lee, Y-D Lee, J. Mater Sci: Mater. Med, 17, 1411–1420 (2006). 16. T. M. Nair, M. G. Kumaran and G. Unnikrishnan, J. Appl. Polym Sci., 93 2606–2621 (2004). 17. S. C. George, K. N. Ninan, G. Groeninckx and S. Thomas, J. Appl. Polym. Sci., 78, 1280–1303 (2000). 18. U. N. Okwu and F. E. Okieimen, Europ. Polym. J., 37, 2253–2257 (2001). 19. C. E. Scott and C. W. Macosko, Polym. Bull., 26, 341 (1991). 20. C. E. Scott and C.W. Macosko, Polymer, 35, 5422 (1994). 21. B. D. Favis, J. Appl. Polym. Sci., 39, 285 (1990). 22. D. Bourry and B. D. Favis, Polymer, 39, 1851 (1998). 23. R. M. Ho, C. H. Wu, A. C. Su, Polym. Eng. and Sci., 30, 511–518 (1990). 24. R. Scaffaro, L. Botta, F. P. La Mantia, P. Magagnini, D. Acierno, M. Gleria, R. Bertani, Polymer Degradation and Stability, 90, 234–243 (2005). 25. M. Mari´c, C. W. Macosko, Polymer Engineering & Science, 41, 118–130 (2001). 26. C. R. Hossein Z. Hot–melt extrusion technique: A review. Iran J Pharm Res., 3, 3–16 (2004). 27. S. Tyagi, A. K. Ghosh, Polym. Eng. and Sci., 42, 1309–1321 (2002). 28. H. M. da Costa, V. D. Ramos, M. C. G. Rocha, Polymer Testing, 25, 498–503 (2006). 29. X. D. Lin, W. L. Cheung, J.Mat. Processing Technology, 63, 476–480 (1997). 30. B. Yin, Y. Zhao, R.-Z. Yu, H.-N. An, M.-B. Yang, Polym. Eng. and Sci., 47, 14–25, (2007). 31. S. C. George, G. Groeninckx, K. N. Ninan and S. Thomas, J. Polym. Sci. Part B: Polym. Phys., Vol. 38, 2136–2153 (2000). 32. Z. Ommen, S. R. Zachariah, S. Thomas, I. Aravind, G. Groeninckx, J. Macromolec. Sci.. Phy., 43, 1025–1043 (2004). 33. J. K. Lee, C. D. Han, Polymer, 41, 1799–1815 (2000). 34. P. Charoensirisomboon, M. Weber, Polymer, 42, 7009–7016 (2001). 35. A. Ning-Jian, W. Qi, J. Mat. Sci. let. 21, 1315–1317 (2002). 36. E. V. Prut, Polym. Sci, Series D, Glues and Sealing Materials, 2, 1–7 (2009). 37. M. D. Elkovitch, D. L. Tomasko, L. J. Lee, Polym. Eng. and Sci., 39, 2075–2084 (1999). 38. Q. Xu, Y. Chang, J. He, B. Han and Y. Liu, Polymer, 44, 5449–5454 (2003). 39. M. Temtem, T. Casimiro, J. F. Mano and A Aguiar-Ricardo, J. Supercritical Fluids, 43, 542–548 (2008). 40. W. Lihotzky-Vaupel, Preparation of moldings by extrusion of blends of natural polymers and other natural materials.(Germany) EP:1321279, 2003. 41. M. D. Elkovitch, L. J. Lee, D. L. Tomasko. Effect of supercritical carbon dioxide on morphology development during polymer blending. Polym Eng Sci., 40, 1850–61 (2000). 42. M. D. Elkovitch, L. J. Lee, D. L. Tomasko. Effect of supercritical carbon dioxide on PMMA/Rubber and Polystyrene/Rubber blending: Viscosity ratio and phase inversion, Polym Eng Sci, 41, 2108–25 (2001). 43. M. S. Sreekala, M.G. Kumaran, R. Joseph, S. Thomas, Comp.Sci. & Technology, 61, 1175–1188 (2001). 44. L. A. Pothen, S. Thomas, G. Groeninckx, Composites, Part A: Applied Science & Manufacturing, 37, 1260–1269 (2006). 45. http://www.azom.com/details.asp?ArticleID=401 46. H. Hout, B. Jensen. J. Polymer Composites, 29, 906–914 (2008). 47. R. Zulkifli, C. H. Azhari, M. J. Ghazali, A. R. Ismail, A. B. Sulong, European J. Sci. Res. 27, 454–462, 1450–216 (2009). 48. S. S. Tavares, N. Caillet-Bois, V. Michaud, J.-A. E. M˚anson, Comp. Sci. and Tech., 70, 797–803 (2010). 49. http://www.ctihuatai.com/index.files/rtmstatic.gif 50. P. A. Sreekumar, R. Saiah, J. M. Saiter, N. Leblanc, K. Joseph, G. Unnikrishnan, S. Thomas, Composite Interfaces, 15, 263–279 (2008). 51. Y. S. Song, J. R. Youn, Polymer Composites, 29, 390–395 (2008). 52. M. J. Tari, A. Balst, J. Park, M. Y. Lin and H. T. Hahn, Composires Purr., A 29A, 65, l–661 ( 1998). 53. G.-W. Leea, N.-J. Leea, J. Janga, K.-J. Leea, J.-D. Namb, Compo. Sci. and Techno., 62, 9–16 (2002).

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˚ om, Composites Part A: Applied Science and Manufacturing, 29, 585–593 (1998). 54. A. Carlsson and B. T. Astr¨ 55. I. Angelov, S. Wiedmer, M. Evstatiev, K. Friedrich and G. Mennig, Composites Part A: Applied Science and Manufacturing, 38, 1431–1438 (2007). 56. C. M. Branco, J. M. Ferreira, P. Faelt and M. O. W. Richardson, Int.J.Fatigue, Vol. 18, 255–263 (1995). 57. M. Kazanci, D. Cohn, Marom, J. Mate. Sci., 36, 2845–2850 (2001). 58. N. O. Cabrera, B. Alcock, E. T. J. Klompen, T. Peijs, Appl Compos Mater., 15, 27–45 (2008). 59. http://www.romeorim.com/process01.shtml 60. J. J. George, A. K. Bhowmick, J. Mater Sci., 43, 702–708 (2008). 61. J. J. George, A. K. Bhowmick, Nanoscale Res Lett., 3, 508–515 (2008). 62. Y. Chen, S. Zhou, H. Yang, G. Gu, L. Wu, J. Colloid and Interface Science, 279, 370–378 (2004). 63. Y. Liu, K. Wang, H. Li, H. Chen and Y. Sun, Rare Metals, 25, 63–67 (2006). 64. Z. Yu, J. Yin, S. Yan Y. Xie, J. Ma, X. Chen, Polymer, 48, 6439–6447 2007). 65. K. Chrissafis, K. M. Paraskevopoulos, S. Y. Stavrev, A. Docoslis, A. Vassiliou, D. N. Bikiaris, Thermochimica Acta., 465, 6–17 (2007). 66. T. Norisuye, Y. Kida, N. Masui, Qui Tran-Cong-Miyata, Y. Maekawa, M. Yoshida, M. Shibayama, Macromolecules, 36, 6202 (2003). 67. http://www.nanotech-now.com/news.cgi?story_id=34180 68. T. Hoare, R. Pelton, Macromolecules, 37, 2544–2550 (2004). 69. Xiaomei Ma and Yacheng Xing, Polymer Bulletin, 57, 207–217 (2006) 70. S. Cardea, A. Gugliuzza, M. Sessa, M. C. Aceto, E. Drioli, and E. Reverchon, ACS Appl. Mater. Interfaces, 1, 171–180 (2009). 71. M. A. Meador, G. L. Gould, LEW Number, LEW–17,685, Research & Technology, NASA Glen Research Centre, 133, (2007). 72. T. T. S. d Mes and S. J. Kenneth, A. Douglas, National Meeting of American Chemical Society (ACS) Chicago, August 20–25 Conf 950801-2 (1995). 73. H.-P. Hentze and E. W. Kaler, Chem. Mater., 15, 708–713 (2003). 74. K. Haraguchi and H-J Li, Macromolecules, 39, 1898–1905 (2006). 75. G. Fleury, G. Schlatter, C. Brochon, C. Travelet, A. Lapp, P. Lindner, and G. Hadziioannou, Macromolecules, 40, 535–543 (2007). 76. L. Zhang , S. Zhang, Electrochimica Acta, 54, 606–610 (2008). 77. K. Haraguchi, T. Takehisa, S. Fan, Macromolecules, 35, 10162–10171 (2002). 78. K. Haraguchi, and L. Song, Macromolecules, 7, 40, 5526–5536 (2007). 79. K. Haraguchi, H.-Jun Li, N. Okumura, Macromolecules, 40, 2299–2302 (2007). 80. S.-D. Naomi, J. Calfred, N. T. Gregory, Biomacromolecules, 9, 2784–2791 (2008). 81. L. F. Gudeman, N. A. Peppas, J. Memb. Sci., 107, 239–248 (1995). 82. T. Nakajima, H. Furukawa, Y. Tanaka, T. Kurokawa, Y. Osada, J. P. Gong, Macromolecules, 42, 2184–2189 (2009). 83. R. Vendamme, S.-Y. Onoue, A. Nakao, T. Kunitake, Nature Materials, 5, 494–501 (2006). 84. L. H. Sperling, An overview of interpenetrating networks, Polymeric Materials Encyclopedia, 5th Ed., J. C. Salamone (Ed.) CRC Press: Boca Raton, FL (1996). 85. S. Rajvaidya, R. Bajpai, A. K. Bajpai, J. Macromolecular Sci, Part A, 42, 1271–1285 (2005). 86. A. P. Mathew, S. Thomas, Ph.D thesis, M.G. University. 87. E. Adem, G. Burillo, E. Bucio, C. Maga˜na, M. Avalos-Borja, Radi. Phy. and Chemi., 78, 549–552 (2009). 88. T. Dongyan, L. Hong, C. Weimin, Ferroelectrics, 265, 259–264 (2001). 89. S.-S. Yoon, J.-H. Kim, S.-C. Kim, Polymer Bulletin, 53, 339–347 (2005). 90. J. S. Lee, J. H. Shin, B. K. Kim, Y. S. Kang, Colloid & Polymer Science, 279, 959–965 (2001). 91. J. D. Zhang, M. J. Yang, Y. R. Zhu, H. Yang, Polym. International, 55, 951–960 (2006). 92. S. Li, W. Zeng, J. Appl Polym Sci., 84, 2347–2351 (2002). 93. O. Grigoryeva, O. Slisenko, A. Bismarck, L. Sergeeva, Macromolecular Symposia, 254, 233–239 (2007). 94. A. S. Vatalisa, C. G. Delidesa, G. Georgoussisb, A. Kyritsisb, O. P. Grigorievac, L. M. Sergeevac, A. A. Brovkoc, O. N. Zimichc, V. I. Shtompelc, E. Neagud, P. Pissis, Thermochimica Acta, 371, 87–93 (2001).

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95. A. Y. Kwok, G. G. Qiao, D. H. Solomon, Chem. Mater., 16 (26), 5650–5658 (2004). 96. M. A. Garce, S. Lara, R. Antoniocruz, A. M. M. Martinez, T. L. Ramirez, A. Beatriz, M. Cepeda Chem. Eng. Comm., 196, 1227–1236 (2009). 97. T. T. Alekseeva, Y. S. Lipatov, N. V. Babkina, N. V. Yarovaya, L. A. Sorochinskaya, Mole. Cryst. and Liquid Crys., 483, 191–204 (2008). 98. R. S. Tıglı, M. G¨um¨us¸derelio˘glu, J Mater Sci: Mater Med., 20, 699–709 (2009). 99. S. Mishra, R. Bajpai, R. Katare, A. K. Bajpai, Microsyst Technol., 14, 193–198 (2008). 100. I. Gitsov and C. Zhu, J. Am. Chem. Soc., 125 (37), 11228–11234 (2003). 101. R. M. Alcˆantara, A. T. N. Pires, G. G. D. Barros, L. A. Belfiore, J. Appl. Polym. Sci., 89, 1858–1868 (2003). 102. H. Chen, X. Deng, X. Hou, R. Luo, B. Liu, J. Macromolecular Sci. R Part A: Pure and Appl. Chem., 46, 83–89 (2009).

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6 Macro, Micro and Nanostructured Morphologies of Multiphase Polymer Systems Han-Xiong Huang Laboratory for Micro Molding and Polymer Rheology, South China University of Technology, Guangzhou, CHINA

6.1 Introduction 6.1.1

Polymer Blends

Physically blending two or more immiscible polymers has become an effective and important route to achieve new high-performance polymeric materials. The blending technique can serve several different purposes. The main purpose is to achieve and tailor a useful combination of a wide range of physical, mechanical, and other properties to the needs of particular end-use applications. Often synergistic effects can be obtained. Other purposes can include improved processability or lower cost. Therefore, polymer blends are used extensively in different industrial applications such as automobile, aeronautic, and packaging. Scientists working in polymers, rheology, and fluid mechanics have been attracted to reveal the phase morphology development mechanisms in polymer blending [1–20]. It is well known that the properties of immiscible polymer blends depend strongly on their ultimate phase morphology represented by the size, the size distribution, spatial arrangement, and orientation of the dispersed domains [21]. The morphology of an immiscible binary polymer blend may be classified as dispersed and cocontinuous. For the former, several categories may be identified, e.g., droplet, fiber, and lamellar dispersion in the matrix [22]. The droplet-matrix morphology presents good stiffness, the fiber-matrix morphology provides great enhancements in unidirectional strength, the lamellar morphology offers excellent barrier properties, and the co-continuous morphology usually exhibits good electrical conductivity, toughness, and stiffness [23, 24]. Even for the droplet-matrix morphology, very slight changes in the average size and size

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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distribution of the droplet can result in obviously different mechanical properties [25, 26] or a large shift in the brittle–tough transition [27] of polymer blends. Therefore, it is generally desired to produce blends with well-defined, stable, and reproducible morphologies [28]. Morphology development is the evolution of the blend morphology from pellet- or powder-sized particles to the micron- or submicron-sized droplets that exist in the final blend. Moreover, during the blending process, an individual droplet at the microscopic level deforms to various shapes, breaks up or interacts with neighboring droplets, and experiences a complicated flow history [29]. Accurate prediction of the blend morphology development in a time scale (batch mixers) or in a length scale (extruders) is thus helpful to prepare polymer blends with specific properties and to design processing equipment [30, 31]. However, polymer blending is complicated because it involves time- and temperature-dependent non-Newtonian materials, and moreover, complex flow field exists in the processing equipment. Therefore, it is difficult to predict how the final morphology is developed [32]. In the pursuit of obtaining new polymer blends, attention has been drawn to systems having more than two phases [33–35]. For ternary immiscible blends, the range of possible morphologies is more complex since three different interfaces can coexist [36]. There are three types of phase morphology [33–35, 37, 38]: (1) two minor components disperse separately in the matrix; (2) one minor component encapsulates the other with a core–shell or composite droplet morphology; and (3) mixed phases of two minor components are formed without any ordered organization. Recently, there is a growing interest in achieving nanostructured blends (nanoblends) of thermoplastic polymers [39–46] because they have the potential for enhanced properties (e.g., transparency, mechanical properties, heat resistance, ductility, creep resistance, and so on) in comparison with conventional immiscible polymer blends with average dispersed-phase sizes on the micron scale. In nanoblends, the dispersedphase domains exhibit length scales of order 100 nm or less. In most previous studies, nanoblends were prepared via in situ polymerization of one monomer in the presence of another polymer or by reactive blending [40–44]. For example, co-continuous nanostructured blends with improved mechanical properties were prepared by reactive melt blending [41, 44]. Using melt-state high-shear (with shear rates exceeding 1000 s−1 ) processing, Shimizu et al. [45] successfully prepared stable nanostructured PVDF/PA11 blends, in which the PA11 domains with a size of several tens of nanometers are dispersed in the PVDF phase. The nanostructured blends can be produced with a wide range of compositions without the addition of block copolymer. Tao et al. [46] demonstrated a continuous, industrially scalable, mechanical process for achieving nanostructured polymer blends. This approach is based on solid-state shear pulverization (SSSP), which yields a fine dispersion of a minor phase in a major phase due to repeated fragmentation and fusion steps within the pulverizer. This solid-state process eliminates certain limitations to the production of nanoblends by melt-state processes, e.g., interfacial tension effects leading to coarsening or coalescence. A PS/PMMA (80/20 wt) blend processed by SSSP and consolidated by platen pressing, without melt processing, exhibits a quasi-nanostructured morphology with many irregular, minor-phase domain with length scales of ∼100 nm or less. When the pulverized blend is subjected to short-residence-time single-screw extrusion, the resulting nanoblend exhibits spherical dispersed-phase domains with a number-average diameter of 155 nm. Thus, SSSP followed by certain melt-processing operations can yield nanoblends. Moreover, nanostructured thermosets have also received attention as their thermal, dielectric, diffusion, and mechanical properties can be systematically tuned depending on the components incorporated [47, 48]. Nanostructured thermosets can be prepared by reaction-induced phase separation [48]. 6.1.2

Polymer and Polymer Blend Nanocomposites

In polymeric nanocomposites, at least one dimension of the dispersed particles is in the nanometer range. Depending on the number of dimensions of the dispersed particles in the nanometer range, three types of

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nanocomposites can be distinguished [49, 50]: (1) isodimensional nanoparticles, such as spherical silica nanoparticles. Their three dimensions are in the order of nanometers; (2) nanotubes or whiskers, in which two dimensions are in the nanometer scale whereas the third is larger, thus forming an elongated structure; (3) nanoparticles with only one dimension in the nanometer range, which are present in the form of sheets of one to a few nanometers thick and hundreds to thousands nanometers long. Clays and layered silicates belong to this type and the composites are known as polymer–clay nanocomposites (PCNs) or polymer-layered silicate nanocomposites (PLSNs). When polymeric materials possess multicomponent phase-separated microstructures at the nanoscale, they can exhibit significantly improved mechanical properties [51]. The nanodomains can constrain the polymer chains or enhance the toughening efficiency of the polymer depending on the type of particle used. However, the extent of improvement is determined by the microstructure, that is, the size, shape, and homogeneity of the particle in the polymer matrix. Of all the potential nanocomposite precursors, those based on clay and layered silicates, which were reported in the patent literature as early as 1950 [52], have attracted the greatest interest, because of the ability of the silicate particles to disperse into individual platelets and the ability to fine tune their surface chemistry through ion exchange reactions with organic and inorganic cations [53]. Numerous benefits have been reported by the addition of clay at low loadings to a variety of matrix polymers. The property improvements can generally include mechanical properties (e.g. tensile strength and modulus), barrier properties, heat resistance, flame retardation, and dimensional stability [54, 55]. The potential property improvements are mainly due to the high aspect ratio of the clay platelets and usually depend on the degree of their exfoliation and dispersion [55]. Three main types of layered silicates exist in polymer matrices [49, 55–60], as illustrated in Figure 6.1. When the polymer is unable to intercalate (or penetrate) between the silicate platelets, an agglomerated microcomposite or tactoid is obtained. In an intercalated structure, a small amount of extended polymer chain

+ Clay Polymer

MELT BLENDING

Tactoid

Intercalated

Intercalated Disordered

Delaminated or Exfoliated

Figure 6.1 Schematic illustration of three main types of layered silicates in polymer matrix. Reprinted from [55]. Copyright (2001) with permission from Elsevier.

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is penetrated between the silicate platelets, but causes less than 2–3 nm separation between the platelets. When the polymer chain further separates the silicate platelets, e.g. by 8–10 nm or more, a delaminated or exfoliated structure may be obtained. Since an ideally exfoliated nanocomposite possesses the largest surface area to volume ratio for the clay, exfoliation leads to better phase homogeneity and the greatest improvements in mechanical and thermal properties. However, complete exfoliation is practically unrealizable, and the best dispersed systems typically consist of a mixture of intercalated and exfoliated structures. Recently, there has been increasing study on the polymer blend nanocomposites. Some works are focused on either improving mechanical properties of polymer blends reinforced with nanoclays [61, 62], or improving impact strength of nanocomposites toughened with a proper functionalized elastomer [53, 63–65]. Some works concentrate on the possibility of using organoclay as a compatibilizer and as a nanofiller for immiscible polymer blends at the same time [61, 62, 66–75]. Since the possibility of direct melt blending via extrusion was demonstrated by Giannelis et al. [76], this method is by far the most common and industrially practical strategy for preparing polymer nanocomposites without in situ intercalative polymerization [54]. Over the last decade, there have been many studies reported on the preparation of polymer nanocomposites [55, 58, 60, 77–81] and polymer blend nanocomposites [65, 66, 72, 82–84] by melt blending. Morphology of polymer blends, polymer nanocomposites, and polymer blend nanocomposites prepared by melt blending can be controlled by many factors, which can be classified as two types, that is, materialand processing-relevant factors.

6.2

Morphology Development Mechanisms of Multiphase Polymer Systems During Processing

Polymer compounding normally starts from solid pellets or powders. It is generally believed that there are three main stages involved in the phase morphology development process of polymer blends during processing: melting of components, deformation and breakup of the dispersed phase, and possible coalescence of the dispersed phase. The former one is the solid–melt transition stage where at least one of the components in the blends is still undergoing melting, that is, the initial stage of the blending process; in the latter two stages all components are totally molten. The mechanism of initial morphology development of blends is crucial in understanding their final morphology, the dynamic equilibrium between the breakup and coalescence of the dispersed phase generally governs the steady-state morphology of blends. Most previous studies on polymer blending have been focused on elementary stages such as the deformation, breakup, and coalescence of the droplet, whereas relatively few studies have been undertaken on the mechanism of initial morphology development. The key objective in preparing polymer clay nanocomposites is to achieve intercalation, and even exfoliation or delamination of the large stacks of silicate nanoplatelets (generally montmorillonite) into single layers or tactoids of small number of layers. Several mechanisms have been proposed for the intercalation, exfoliation, and dispersion process of organoclay in polymers during melt-mixing. 6.2.1

Initial Morphology Development in Polymer Blending

A number of investigations on the effect of mixing time on the blend morphology demonstrated that under normal processing conditions, rapid morphology development usually occurs within the initial 2–5 min of the mixing in an internal mixer [30, 85–88], or in the initial melting zone (1–3 times of screw diameter) of a twin-screw extruder (TSE) [89, 90] or single-screw extruder (SSE) [91], which is in conjunction with the melting of the components. The dispersed phase size changes less after the polymers are completely melted.

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The research by Burch and Scott [92] also confirmed the importance of the softening/melting regime to the morphology development in low-viscosity-ratios (η = 1.12, 0.11, 0.03, 0.003, and 0.00015), miscible blends of SAN/PMMA and LDPE/LDPE. For PA66/PEG immiscible blends with an extremely low viscosity ratios (ηr = 3–4 × 10−5 ), experiments carried out with two mixing time (5 and 10 min) by Leblanc et al. [93] showed no significant difference of dispersed PEG phase size distribution. For HDPE/PS/PMMA ternary blend, Reignier and Favis [94] found that the composite droplet morphology occurs within the first 2 minutes of mixing and remains stable thereafter. 6.2.1.1

Surface Erosion Mechanism

In 1990, Plochocki et al. [95] carried out the investigation on initial morphology development during polymer blending. They compounded PS with LDPE in a variety of industrial mixers. They tentatively proposed that an ‘abrasion’ or ‘erosion’ mechanism is responsible for the initial stage of the dispersion process, in which solid or only partially softened pellets are abrased against the walls of processing equipment. They suggested that there is abrasion (similar to rubber abrasion) between the two types of pellets and between the pellets and wall of processing equipment. Subsequently, Potente et al. [90] carried out the investigations on the mechanisms that take place in the initial formation and further development of the morphology in the melting section of a co-rotating TSE for the immiscible PP/PA6 blend and also found the surface erosion mechanism. The obtained typical scanning electron microscope (SEM) micrographs (Figure 6.2) clearly show that smaller volumes of the dispersed PA6 phase that melt on the surface of the individual granules become detached from the melting granule surface due to the prevailing flow conditions and are rapidly converted into small droplets (<10 μm diameter), which display dimensions of the same order of magnitude (μm range) as those observed at the end of the extruder. Moreover, an evaluation of SEM micrographs (Figure 6.2) shows that, in the melting section of the TSE, virtually all the breakup mechanisms that can essentially be distinguished, such as quasi-steady drop breakup, folding, end pinching, and breakup due to capillary instabilities and coalescence processes, take place in parallel. Using on-line visualization, Sundararaj and co-workers studied the deformation and breakup of PC pellets in a PE matrix under complex flow during blending process of the PE/PC blend in a TSE [16] and the deformation and breakup of a single PC drop in a PE matrix under simple shear flow for 6 < ηr < 60 [12], respectively. The results showed that a major mechanism is erosion from the surface of the mother drop in the form of thin ribbons and streams of small droplets. Also using flow visualization, Mighri and Huneault [17, 96] investigated the deformation and breakup under simple shear of single molten polymer drop (PS or EPR) in a polymer matrix (PE or PP). At high shear rates, the most striking non-Newtonian effects are the surface erosion and the drop splitting mechanisms. The authors suggested that only in molten polymer blends, typically several orders of magnitude more viscous than Newtonian ones, can the shear stresses at the interface be sufficient for particle erosion to occur. The erosion of the drop phase occurred by the local stretching of small and thin domains on the surface causes their rupture from the rest of the drop by tensile failure. The drop splitting occurred after the elongation of the droplet parallel to the vorticity axis because of elastic normal forces generates along streamlines inside the drop. This splitting is related to a rocking motion that brings the end of the fluid fiber in different velocity layers, resulting in its quick tearing apart. The particles eroded off the main droplet surface are very fine, in the range of 10–50 μm, and result in a significant reduction in main drop size before its final breakup. Actually the erosion breakup phenomenon has already been studied in many other fields, such as agglomerate dispersion [97], drug delivery, and rock erosion. The erosion mechanism observed in the aforementioned experiments can be explained as follows. Chen et al. [98] numerically simulated the deformation and breakup mechanism of PC or PE drop in a PE melt under shear flow. When the shear rate (γ˙ ) is increased, both shear

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Figure 6.2 Surface erosion mechanism in initial morphology development in melting section of a TSE for PP/PA6 blend. Reprinted from [90]. Copyright (2001) with permission from John Wiley and Sons.

stress and normal stress are increased because the former is proportional to γ˙ and the latter is proportional to γ˙ 2 . A competition between both stresses results in more irregularity of drop shape. Their simulation showed that the shear rate and shear stress at the interface are much higher than those either inside the dispersed phase or inside the matrix phase. This results in that the drop surface is much less viscous than in the center of the drop and allows material to be easily peeled off of the drop surface. Therefore, when the shear rate reaches a critical value, the drop can no longer sustain the material at the surface, and surface erosion begins as the drop releases ribbons and droplets into the matrix. 6.2.1.2

Sheet Formation Mechanism

As early as 1985, Min and White [99] studied the flow behavior of elastomers and molten plastics in a Banbury-type mixer. They observed the stretching, tearing, ‘sheeting out’, and banding in elastomers, and ‘sheeting out’ in plastics. The ‘sheeting out’ phenomenon was observed as the colored sample sheet is sheared as a thin film layer on the chamber wall. More systemic studies on the initial stage during blending using batch mixers [30, 85, 86, 100], TSEs [89, 100], and miniature mixers [100] were performed by Macosko and co-workers for several polymer blends. Their results showed that the initial stage of blending can be described by a ‘sheeting’ breakup mechanism, as shown in Figure 6.3. During mixing, the molten components are stretched and folded into thin sheets or ribbons due to the complex flow field, reaching a critical thickness when capillary wave instabilities lead to

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Holes form in the ribbon due to interfacial instability.

167

The lace breaks into irregularly shaped pleces with diameters approximately equal to the ultimate particle size.

2 μm 10 μm

V

Sheet or ribbon of the dispersed phase forms in the matrix.

10 μm

The size and concentration of holes becomes sufficient for a lace structure to be formed.

2 μm

Drop and cylinder breakup form the final spherical particles.

Figure 6.3 Schematic of sheeting mechanism for initial morphology development in polymer blending. Reprinted from [30]. Copyright (1995) with permission from Elsevier.

formation of ‘holes’ and a two-dimensional lacelike structure, which is stabilized by the antitropic network of the other component that coalesces between the holes. During further mixing, the lacelike structure becomes unstable due to the effects of flow and interfacial tension and is broken into irregularly shaped particles or into cylinders and eventually spherical droplets via Rayleigh-type instabilities. This mechanism results in fast formation of small dispersed droplets, which are nearly the same size as those observed at long mixing times. Continued mixing primarily reduces the size of the largest droplets in the size distribution. So this proposed mechanism results in the generation of very small droplets at very short timescales. The aforementioned sheet formation was also observed during blending in a SSE by Lindt and Ghosh [91], who investigated the blending of PS with styrene-butadiene copolymer and with EVA, respectively. Through drawing the screw and examining the frozen ribbon peeled off the screw channel, they observed that the melting pellets produce fine lamellar structures (thickness of the order of μm) that extend over much of the melting zone and then are broken up by capillary forces. The mention above shows that during the initial stage of polymer blending, lamellar structures or sheets are formed [30, 85, 89, 91, 101] and morphology development via sheet breakup is an effective way to achieve quick reduction in particle dimension [101, 102]. This can be explained as follows. For polymer blends, in the initial stage of breakup, the interfacial tension is negligible when compared with the normal stresses [11]; therefore, it is much easier to form a sheet instead of the cylinder that is usually formed in Newtonian systems. Willemse et al. [102] theoretically explained the rapid decrease of the sheet thickness under shear flow by using simplified equations. The final phase dimensions are found to be largely determined by the sheet thickness at the onset of breakup of the sheets.

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M

T

P

Melt

Figure 6.4 Schematic of morphology development of an immiscible blend starting from solid pellets. Reprinted from [21]. Copyright (2001) with permission from John Wiley & Sons.

Li and Hu [21] described the morphology development of the PP/PA6 (80/20) blend starting from solid pellets as in Figure 6.4. A solid pellet of PA6 (S) is first transformed to the melt (M). Usually, the initial molten layers stripped from the solid pellet have much higher capillary numbers than their critical capillary numbers. Therefore, they undergo transient affine deformation to stretch to thin sheets or threads (T), which break up into small particles (P) due to interfacial instability. Moreover, the authors suggested that if the polymer forming the minor phase has an higher melting/softening temperature than the one forming the matrix, the most favorable conditions for obtaining a fine morphology appear to be the following: the rate of melting of the dispersed phase < the rate of deformation and breakup of the dispersed phase melt into small droplets < the rate of stabilization of these small droplets using an adequate compatibilizer.

6.2.2 6.2.2.1

Deformation and Breakup of Droplet Approaches and Main Findings of Flow-induced Droplet Deformation and Breakup Studies

Polymer blending is a complex time-dependent process which involves flow-induced breakup and coalescence of viscoelastic fluids [14, 101, 103–105]. In detail, the early mixing stage is controlled by breakup, and the final stage is dominated by the balance between drop breakup and coalescence [103, 105]. A fundamental understanding of deformation and breakup mechanisms of isolated polymer droplets under a well controlled-flow field is helpful to a better control over blend morphology development [6, 17]. Therefore, flow-induced droplet deformation has long fascinated fluid dynamicists. Droplet deformation and break experiments are generally performed in optical flow cells with shear or elongational flow. Shear flow field is commonly generated by two devices: one is parallel plates and the other is Couette flow cells. There may be four types of parallel plate devices: two transparent plates, which can be either counter-rotating or only one of them may rotate [4, 6, 13, 101, 106–109]; cone-and-plate [9, 110, 111]; sliding glass plate devices in combination with a movable microscope [112–114]; parallel-band cell, which uses two parallel bands moving in opposite directions [115, 116]. Parallel plates can be used to study very viscous polymers. However, the flow field generated by rotating the parallel plates is not homogenous [14]. In Couette flow cells, two concentrical cylinders can be either counter-rotating or only one of them may rotate [8, 11, 12, 117]. Elongational flow field can be generated by using a converging channel [118–121] or a four-roll mill [122–126]. Moreover, an eccentric cylinder device can be used to create chaotic flow [127–131] or complex flow [132] for studying the droplet deformation. Two experimental approaches have been used to study the behavior of isolated droplets under well-defined flow fields. The one is to investigate the steady-state shapes of deformed droplets and the critical conditions required for breaking the droplets under steady shear flow [5, 6, 111, 112, 115, 133–135]; the other is to use step-strain experiments, wherein the effects of material and flow parameters on the transient deformation and shape relaxation behavior of isolated droplets can be separately studied [111]. Taylor [133, 115] performed the pioneering work on the steady-state deformation and breakup of a single Newtonian drop suspended in a Newtonian matrix. By balancing the interfacial stress (σ ) and the shear

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stress, he put forward a relationship to predict the largest droplet diameter (D) that would be stable for small deformations in Newtonian fluids: D=

19ηr + 16 Ca 16ηr + 19

(6.1)

where Ca is the capillary number and expressed by: Ca =

ηm γ˙ 2σ/D

(6.2)

where ηm is the viscosity of matrix phase. That is, the Ca is the ratio of the two opposing forces, i.e. the viscous stress that tend to deform the droplet to the counteracting interfacial stress that tend to resist the deformation and keep the droplet shape spherical. Taylor’s theory means that droplet deformation is controlled by two dimensionless parameters, i.e. the Ca and the ηr . Moreover, this theory is valid for all small deformations in Newtonian fluids when ηr < 2.5. Guido and Villone [112] demonstrated that the steady-state three-dimensional shape of a droplet under simple shear flow can be well described by an ellipsoid having three different principal axes (shown in Figure 6.5), in which the steady-state length of the short axis of the ellipsoid measured in the perpendicular plane (plane containing the flow direction and the vorticity axis) is slightly larger than that measured in the shear plane. Taylor’s theory was later verified for Newtonian fluids by Karam and Bellinger [136] and by Grace [134]. Many other experimental observations, asymptotic analyses, and numerical simulations have been devoted

θ

b/2 a/2

c/2

y

x z

Figure 6.5 Schematic of an isolated droplet undergoing simple shear flow between two parallel plates. Reprinted from [112]. Copyright (1998) The Society of Rheology.

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to understand the drop behavior under both small and large strains and to assess the conditions of drop breakup under various flow geometries and conditions. Extensive literature exists concerning this research topic, which dates back to the original contribution by Taylor [133, 115], and spans to our days. The main findings of such studies can be summarized as follows. When the Ca is small, the initial spherical droplet deforms gradually and reaches an equilibrium ellipsoidal shape with the major axis of the ellipsoid oriented at an angle close to 45◦ with respect to the flow direction. When the Ca is increased, the droplet deformation increases and the deformed droplet axis shifts toward the flow direction, that is, the angle is smaller than 45◦ . If the Ca reaches the critical capillary number (Cacrit ), the droplet transforms from an ellipsoidal shape to a sigmoidal shape in which the central part of the sigmoid stretches and simultaneously becomes thinner with applied strain. Finally the droplet breaks up into smaller droplets when the Ca goes beyond the Cacrit . Empirical approaches were used to correlate the final droplet size to the known parameters such as initial droplet size [134] or maximum stable droplet size [137]. Bigio et al. [137] found that the maximum daughter droplet size is approximately the same as the maximum stable droplet size at the corresponding shear rate, and the mean daughter droplet size is about 70% of the maximum stable droplet size. Cristini et al. [138] showed that the droplet size distribution in simple shear flow depends only on the average initial droplet size and the shear rate. Employing the aforementioned step-strain experiments, the deformation and relaxation behavior of isolated droplets can be separately studied. Takahashi and co-workers [139, 140] studied the shape recovery of a droplet in an immiscible polymer matrix for blends with a value of ηr less than 1. The droplet exhibits a flat ellipsoidal shape just after applying a moderate step strain. Subsequently, the droplet relaxes back to a spherical shape through a series of intermediate shapes, i.e., cylindrical, dumbbell, and ellipsoid of revolution. The revolution is driven by the tendency toward reduced interfacial area. The length of the major axis of the ellipsoid is slightly larger than that predicted by affine deformation; however, the orientation angle of this axis with respect to the flow-vorticity plane is in good agreement with that predicted by affine deformation. Guido et al. [141] and Hayashi et al. [142] studied the droplet behavior after imposition of reversing step shear strains and large double-step shear strains, respectively. Some works have been dedicated to the study on the deformation and breakup of the droplet subjected to oscillatory shear and extensional flows, where both the amplitude and frequency of deformation can be carefully controlled. Using in situ microscopic visualization, Wannaborworn et al. [143] carried out a pioneering work to observe the deformation and breakup of immiscible Newtonian droplets with a ηr of unity under oscillatory shear flows in a range of strain amplitudes and frequencies. For moderate strain deformations, it was found experimentally that the deformation parameters oscillate sinusoidally between a maximum and a nonzero minimum value, and the numerical simulation is able to capture both this and start-up effects. For large strain deformations, droplet breakup mechanism occurs only through the end pinching at the point of maximum strain. Cavallo et al. [144] investigated experimentally the deformation of an isolated droplet in an immiscible liquid undergoing oscillatory shear flow as a function of frequency and up to moderate amplitudes. It was found that in the small amplitude range, the time-dependence of the axes of the droplet is still harmonic, but shifted in phase with respect to the applied stress, and its amplitude is linear with strain. At higher strains the deviations from a simple sinusoidal response become more evident. The dynamics of a single Newtonian droplet immersed in a Newtonian matrix (with a ηr of 6) subjected to large amplitude oscillatory shear flow (LAOS) was later studied by Guido et al. [145]. The droplet shape was found to oscillate at higher harmonics of the forcing frequency when the Ca is increased. Renardy [146] conducted numerical simulations for the same fluids and flow properties as the experiments of Guido et al. [145] under LAOS. Computational results on the evolution of droplet length and inclination angle are shown to confirm the experimental finding in Guido et al. [145]. In particular, the computed velocity fields around the droplet are shown to elucidate the over-rotation, which is a mechanism for the experimentally observed harmonics.

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In simulating the deformation of viscous droplet in a planar oscillatory extensional flow, Li and Sarkar [147] showed that the oscillation leads to decreased deformation and bounded droplet shapes for conditions for which steady extension results in droplet breakup. Using an optical shear flow cell, Janpaen et al. [148] studied experimentally the deformation and breakup of the Newtonian and Boger droplet in oscillatory shear flow. The results showed that the deformation amplitude parameters generally decrease with increasing ηr , time scale ratio, and droplet elasticity. The dependences of the deformation amplitude parameters on Ca are generally linear up to a certain value for Newtonian droplets regardless of the ηr and time scale ratio. The dependences become totally nonlinear with increasing droplet elasticity. Droplet viscosity and elasticity generally impede oscillatory deformation. At low time scale ratio, droplet deformation varies with time under quasiequilibrium state; for large time scale ratio, the visually apparent deformation can be referred to be in a frustrated state. When time scale ratio increases, the Cacrit for breakup increases linearly, the number of resultant daughter droplets and the number of cycle required for breakup to occur also increase, and the apparent breakup pattern changes from the dumbbell type to the end-pinching type. Also using optical techniques, Deyrail et al. [149] investigated the deformation and breakup of PDMS droplets embedded in PIB under LAOS at various frequencies and deformation amplitudes up to 800%. It was observed that depending on the deformation amplitude, the shape of the droplet at the flow reversal point varies from a slightly deformed sphere to a highly deformed ellipsoid. On the theoretical side, the finite element method [150, 151], boundary element method (BEM) [118, 119, 152–155], boundary integral method (BIM) [156], volume-of-fluid method [157–160], and front-tracking methods [147, 161, 162], were used or proposed to simulate the droplet deformation. Using the BEM, Khayat and co-workers studied the droplet deformation in shear and elongational flows with relevance to mixing. A boundary element analysis was carried out to simulate the droplet deformation in the screw channel of a TSE [152]. Experiments were carried out in rectangular tubes for droplets initially located on and off the tube axis. Good agreement was obtained upon comparison between simulated and experimental results despite the limiting assumption of two-dimensional analysis [118, 119]. A shear- or elongation-dominated droplet deformation depends on the size of the droplet relative to the channel dimensions, and its position relative to the axis of the channel. The influence of the ηr and interfacial tension on the droplet deformation was also examined. The effects of shear thinning were examined for a droplet moving in a Newtonian ambient fluid [153]. All these works were limited to two-dimensional flows. An adaptive BEM was so proposed for analyzing the three-dimensional droplet deformation in confined flow. The adaptive method is stable as it includes remeshing capabilities of the deforming interface between droplet and suspending fluid, and thus can handle large deformations [154]. Wang and Dimitrakopoulos [155, 163] developed a novel three-dimensional spectral BEM for studying the droplet deformation. Cristini et al. [138, 164, 165] developed a new threedimensional BIM for simulating the droplet breakup process in viscous flows. Their implementation of a fully adaptive surface discretization procedure provides efficient and accurate resolution of the highly deformed droplet shapes. Feigl et al. [150] proposed a three-step numerical procedure to simulate the deformation of droplets in a mixed flow field. Finite element and numerical particle tracking techniques are used to obtain the history of shear and elongation rates along a particle trajectory in the flow field, and from this history, BIM is used to determine the deformation a droplet would experience along this path. The authors then used this method to investigate the deformation and breakup behavior of droplets in the annular gap flow between two eccentric cylinders in which the inner cylinder rotates at a constant angular velocity. It was found that dispersing capability increases with increasing eccentricity. The proposed simulation procedure can serve as an effective tool to determine droplet breakup in dispersing geometries and hence to optimize the dispersing procedures in complex flows. Anderson and co-workers also proposed a new BIM to study the deformation of the confined droplet between two parallel walls for systems with a ηr of unit [166, 167] and non-unit [168].

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6.2.2.2

Types of Droplet Breakup Mechanisms

Four mechanisms have been observed for droplet breakup [10, 169]: capillary breakup [128, 133, 170, 171], necking [128, 172], tip-streaming [171], and end-pinching [111, 143, 148, 173]. Among them, end-pinching is a form of nonuniform breakup wherein fiber disintegration proceeds initially from the ends of the thread and can result in a distribution of particle sizes. It was found that this type of breakup depends on the ηr and initial length-to-diameter ratio of the elongated droplet. Zhao [174], investigating the breakup of Newtonian droplets in dilute Newtonian emulsions in simple shear flow, found that the final droplet size distribution is intimately linked to the droplet breakup mechanism, which depends on the Ca and ηr . For Ca < 2Cacrit , end-pinching is the dominant droplet breakup mechanism. For Ca > 2Cacrit , breakup dynamics are strongly controlled by ηr . At 0.1 < ηr < 1.0, capillary instability is the dominant breakup mechanism, and thread radius and wavelength at breakup are uniform and independent of the initial droplet sizes. The average final droplet size is inversely proportional to the shear rate. At 1.0 < ηr < 3.5, droplets break up via long wavelength capillary instability but large satellite droplets lead to a bimodal droplet size distribution. At ηr < 0.1, droplets break up via a complex re-breaking and collision process. Decreasing the ηr results in more polydisperse emulsions. Lin et al. [15] illustrated four types of polymer droplet breakup mechanisms (as shown in Figure 6.6) occurred in a simple shear flow field. (1) Tip streaming [115, 134, 175–178]: a well-known breakup mechanism even in Newtonian systems. The droplet assumes a pointed sigmoidal shape, and streams of small droplets are released from the tips of a pointed mother droplet along the flow direction. (2) Parallel breakup or sheet breakup [11, 101]: the droplet stretches into a thin sheet along the flow direction and breaks up into two or more smaller daughter droplets, with one or several satellite droplets between them. (3) Surface erosion: surface erosion from the droplet in the form of thin ribbons and streams of small droplets. (4) Vorticity elongation and breakup [4, 8, 27, 96, 107, 108, 179]: the droplet breaks after being elongated in the vorticity direction, that is, perpendicular to the flow direction. In addition to mechanism 1, all other three mechanisms are unique to viscoelastic systems. Mechanisms 2 and 3 also occur in the initial stage of the blending process, see Section 6.2.1. As noted in Section 6.2.1.2, sheeting phenomenon was observed in the initial stage during blending. Sundararaj et al. [101] showed that sheets are also formed when both components are molten. They visualized the deformation and breakup of PP drop inside a PS matrix subjected to steady-state shear between transparent parallel plates. They observed that PP drops are first stretched out into sheets and then broken into cylinders

Tip Streaming

Parallel Breakup Polymer Drop

Erosion

Vorticity Direction

Flow Direction

Vorticity Alignment and Breakup

Figure 6.6 Schematic of four types of polymer droplet breakup mechanisms in simple shear flow. Reprinted from [15]. Copyright (2005) with permission from American Chemical Society.

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Figure 6.7 Two breakup modes of sheet formed during initial shearing. Reprinted from [14]. Copyright (2004) with permission from Elsevier.

via different mechanisms, such as hole formation, growth, and coalescence to give ligaments oriented parallel and perpendicular to the flow direction. Lin and Sundararaj [14] subsequently studied the deformation and breakup of a PC drop sheared inside a PE matrix at low shear rates. Sheet formation is observed during the initial shearing of the drop. Then the sheet is broken up into smaller droplets via either stretching into a thin thread (mode A) or forming a thin cylindrical tip (mode B), as shown in Figure 6.7. In mode A, as time goes on, the sheet is stretched into a thread, which becomes thinner and thinner, and then twists in the matrix (left, Figure 6.7a). Finally, a typical thread breakup occurs, where the thread breaks into many smaller droplets (right, Figure 6.7a). In mode B, a thin cylindrical tip develops projecting out of the elongated sheet (left, Figure 6.7b). The initial drop breaks up into smaller drops, as shown on right in Figure 6.7b, which shows such a smaller droplet with small threads entangled around it. Lin and Sundararaj found that the sheet forms at a critical strain (γ˙ tc ) or critical time (tc ). The strain, γ˙ tc, is related to the ηr , stress ratio (Sr ), and droplet Deborah number (De). The Sr is defined as the ratio between the matrix breakup stress, made up of the matrix normal stress and viscous stress, and the droplet restoring stress, made up of the droplet normal stress and the interfacial stress, which can be simplified as: Sr =

2G m + ηm γ˙ 2G d + 2σ/D0

(6.3)

where D0 is the diameter of the pellet or droplet; G m and G d are the elastic modulus of the matrix and droplet, respectively. Pellet breakup is also controlled by the rate of deformation of the pellet phase and its rate of

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relaxation. If it relaxes quickly in comparison to the process time, then the deformation would be expected to be more like the Newtonian case (where the relaxation time is zero). However, if the pellet phase relaxes slowly in comparison to the process time, the deformation is more elastic. The ratio of the relaxation time of the material to the characteristic process time is expressed as the De: De = λγ˙

(6.4)

where λ is the characteristic relaxation time of the pellet phase. The Sr is used to characterize the sheet formation during the droplet deformation and breakup process. It was found that the ηr , Sr and De of the system could be used to predict the droplet deformation and breakup. For the systems studied by Lin and Sundararaj [14], the droplet is easier to breakup at lower ηr and higher Sr . When the Sr is less than 4, the De of droplet decreases with increased Sr ; therefore, the normalized critical time for the droplet to form a sheet is fairly constant as Sr increases and the droplet is broken up through the aforementioned mode B. On the other hand, when the Sr is higher than 4, the De increases with increased Sr , resulting in a delay in sheet formation and the droplet is broken up via mode A. Sundararaj et al. [101] presented a map showing the different conditions, where either sheets or threads may be formed using the ratio of normal stresses and the De as parameters. It was shown that sheets can easily be formed in shear flow and that threads, however, are formed only at low shear rates (<5 s−1 ). Sheets are not transformed into threads because the interfacial forces are far too low under most circumstances to change the shape of the cross-section of the sheets significantly. Using on-line visualization, Lin et al. [12] studied the deformation and breakup of a single PC droplet in a PE matrix under simple shear flow for 6 < ηr < 60. The results showed that a major breakup mechanism is erosion from the mother droplet surface in the form of thin ribbons and streams of small droplets. Also using flow visualization, Mighri and Huneault [17, 96] investigated the deformation and breakup under simple shear of single molten polymer droplet (PS or EPR) in a polymer matrix (PE or PP). At high shear rates, the most striking non-Newtonian effects are the surface erosion and the droplet splitting mechanisms. The authors suggested that only in molten polymer blends, typically several orders of magnitude more viscous than Newtonian ones, can the shear stresses at the interface be sufficient for particle erosion to occur. The droplet splitting is related to a rocking motion that brings the end of the fluid fiber in different velocity layers, resulting in its quick tearing apart. The particles eroded off the main droplet surface are very fine, in the range of 10–50 μm, and result in a significant reduction in main droplet size before its final breakup. In investigating the breakup of viscoelastic droplets (Boger fluids) in Newtonian matrix (PDMS) undergoing simple shear flow, Li and Sundararaj [117] observed that large droplets break up in the flow direction, whereas small droplets break up along the vorticity direction. The transition in the breakup mechanism abruptly occurs at a critical droplet size, which is related to viscosities, interfacial tension, the elastic property, and the critical capillary number for its Newtonian counterpart with the same ηr . Moreover, the critical shear rate changes dramatically when going from one breakup mechanism to another. Flow-induced droplet deformation experiments have been widely employed to study the effects of materialand processing-relevant factors on the blend morphology, which will be discussed in corresponding sections. 6.2.3

Coalescence of Droplet

The coalescence of droplets in a molten polymer blend, being an important factor influencing the dispersed phase morphology, can be described as four steps [180, 181]. (1) The approach of the droplets as a result of the flow field or, under quiescent conditions, as the result of Brownian motions. The higher the concentration of the dispersed phase, the closer the droplets and the higher the probability of collision. In practice, it is assumed that this step is only important in blends with a volume fraction of the dispersed phase lower than 10%; for

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higher concentrations, many droplets practically touch each other. (2) The deformation of droplets during collision along with the removal of the matrix film trapped between the colliding droplets (‘film drainage’). This is often considered as the decisive step in coalescence, and is related to interfacial mobility and droplet size. (3) The rupture of the remainder of the matrix film, at the thinnest spot, usually by the formation of a hole. (4) The coalescing droplets merge to form again one single droplet. Ziegler and Wolf [182] found that the droplet size distribution becomes bimodal for low shear rates and volume fractions of the dispersed phase during the early stages of shear induced coalescence. Coalescence of dispersed phases makes difficult to get fine polymer blend morphology. Coarsening caused by coalescence in blends has been studied in a number of works [103, 105, 180, 183–186]. The results showed that parameters affecting the coalescence involve the volume fraction of the dispersed phase, the interfacial tension, the viscosity of each phase, and the phase size. It was observed that a dispersed phase concentration of 0.5–1 wt% can give rise to flow-induced coalescence [103]. The interfacial tension is the ultimate driving force for coalescence. Low interfacial tension of an immiscible blend, such as PVC/LLDPE blend with an interfacial tension of about 3.4 mN/m [187], can inhibit the coalescence over a wide content range. The blend with high interfacial tension, such as PE/PA6 blend (about 13 mN/m), exhibits larger phase sizes and more composition dependence. The matrix viscosity plays a key role in the time required for film drainage between colliding droplets. Most studies showed that increased matrix phase viscosity reduces the probability for coalescence [103]. For HDPE/PS blends, Lyu et al. [188] found that the coalescence rate increases first and then decreases with increasing the ηr of the blend. There exists a maximum in the coalescence rate within a ηr range between 0.1 to 1. Moreover, it is reported that coalescence efficiency decreases with increasing the droplet size and its difference or the density difference between dispersed phase and matrix [103]. At a given shear rate and constant volume fraction of dispersed phase, the collison frequency under flow-driven coalescence is directly proportional to D−3 (D: droplet size) [189]. To quantify the coarsening of dispersed phases in blends, many studies have been carried out in a quiescent or quasi-quiescent state. Under these conditions, diffusion dependent phenomena were shown to significantly underestimate the rate and extent of dispersed phase coarsening in blends [105] caused by the high viscosities and extremely low mutual solubilities of polymer–polymer pairs. A theoretical description of shear-induced coalescence was first developed by Smoluchowski [190]. Neglecting the droplet interactions and assuming that the droplets are of equal radius R, the concomitant growth of droplet size can be calculated as follows: d[ln(R)]/dγ = φd (4/π )(2 − 22/3 )C

(6.5)

where φ d is the local volume fraction of the dispersed phase, γ is the shear strain, C is the coalescence efficiency. Equation (6.5) shows that the droplet size grows exponentially with γ at constant C. The Smoluchowski theory is focused mostly on the aforementioned first step of coalescence, and predicted a much slower (practically immeasurable) increase in droplet size with time than that obtained experimentally [180]. This is because that the decisive step in flow-driven coalescence is second one aforementioned. This is generally governed by the interfacial mobility, which is quite high for polymer melts. As a result, a relatively high coalescence rate is encountered in polymer blends. Later, Tokita [183] equated the coalescence rate derived by Smoluchowski to an expression for the drop breakup rate during processing, proposing an expression for estimating the equilibrium diameter of the dispersed particles in polymer blends: D≈

24Pr σ π ηm γ˙

  4Pr E D K 2 φd + φd π ηm γ˙

(6.6)

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where Pr is the probability that droplets coalesce after their collision, EDK is the bulk breaking energy. It was suggested that Eq. (6.6) overestimates the influence of shear stress on the particle breakup. Considering that in shear flow, two spherical droplets approach each other at less than the critical separation distance (hc ), Elmendorp and Van Der Vegt [184] defined a critical coalescence time (tc ) as the time that passes between the arrival of a droplet at a liquid/liquid interface and the rupture of the intervening film. In the case of mobile interfaces, the tc is estimated as: tc =

  D 3ηm D ln 4σ 4h c

(6.7)

Taking into account the competing processes of breakup and coalescence of the dispersed phase, Fortelny et al. [191] proposed the following equation for predicting its equilibrium diameter: D=

8σ Pr 2σ Cacrit + φd ηm γ˙ π ηm f 1

(6.8)

where the first term on the right represents the critical droplet diameter calculated from the Cacrit , f 1 is the slope of a function describing the frequency of droplet breakup at Cacrit . This relationship contains several parameters which are difficult to be quantified for the blending of viscoelastic polymers, so its application is somewhat limited. Using PA6/HDPE blends, Gonz´alez-Nu˜nez et al. [192] detected a limiting dispersed phase concentration for coalescence in extensional flow. At low draw rates, the deformation is independent of the dispersed phase concentration. As the draw rate increases, the final state of dispersed phase deformation depends on the composition. At 1–4 vol% of PA6, the average volume of the dispersed phases remains constant (no coalescence effects) during the stretching process. However, at concentrations higher than 5 vol% of PA6, coalescence begins to play a role. These results clearly indicate the limiting dispersed phase concentration for coalescence in this blend under extensional flow. Some works have also been carried out on the coalescence occuring in ternary blends. Reignier and Favis [94] observed the coalescence process for composite droplets in blending HDPE/PS/PMMA system. Composite droplets clearly experience a dual process for static coalescence: first between dispersed PMMA subinclusions within composite droplets, and second between composite droplets themselves. A complete transition from multiple subinclusion particles within a given shell to a core–shell structure was observed upon annealing, that is, at long times. Furthermore, it is clearly shown that coalescence also occurs under dynamic mixing conditions, since the composite droplet size increases with composition. The dynamic coalescence is controlled by the thickness of the PS shell layer. The coalescence behavior of the composite droplet is essentially equal to that of a pure PS dispersed phase above a critical PS shell thickness of about 0.2 μm. A number of researchers have shown that the droplet coalescence can be suppressed upon addition of compatibilizers to blends and several mechanisms can be likely relevant for this coalescence inhibition. Among them, two mechanisms (as illustrated in Figure 6.8) have received the most attention. One is steric repulsion between two approaching droplets (Figure 6.8a) [103]. Block copolymer layers at the interfaces act as squeezed springs when droplets approach in a flow field. In this steric repulsion it is assumed that the block copolymers do not move on the interface, as expected for static coalescence. Steric hindrance becomes more important and coalescence is suppressed more with increasing coverage [193, 194]. Sundararaj and Macosko [103] predicted that the suppression of coalescence is more efficient with higher molecular weight block copolymers. Another mechanism is the Marangoni stress present at the droplet interface in the gap between two approaching droplets (Figure 6.8b) [195, 196]. The Marangoni stress is induced by a gradient of block

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Figure 6.8 Schematic of two mechanisms for block copolymer suppression of coalescence: (a) steric repulsion and (b) Marangoni stress (surface tension gradient). Reprinted from [194]. Copyright (2002) with permission from American Chemical Society.

copolymer concentration at interfaces, which is caused by the squeezing flow between the two approaching droplets (film drainage). As two droplets approach each other, the escaping fluid sets up ‘fountain flows’ inside the droplets, which spread out the velocity gradient and eliminate the singular lubrication force. Since this Marangoni stress acts tangentially on the block copolymer layers that are removed from the gap towards the backside of the droplets and tries to force them back into the gap, the rate of film drainage slows down and can possibly be suppressed completely. The block copolymer concentration necessary for preventing coalescence is higher at higher shear rate. Ha et al. [196] showed with a scaling analysis that the very weak flows typical of a successful coalescence event are sufficient to produce the required surface tension gradients to make the Marangoni contributions to the stress balance significant. Moreover, these two mechanisms might simultaneously contribute to the coalescence suppression. In these conditions, the degree of interface coverage, in relation to the mobility of block copolymers, helps elucidate which mechanism plays the most important role [189]. Separating coalescence from breakup, Lyu et al. [194] studied the coalescence suppression by adding a PS-PE block copolymer into semi-concentrated PS/HDPE blends. Even 0.5% copolymer significantly suppresses the coalescence of HDPE droplets in a PS matrix. A minimum concentration of copolymer is required to suppress coalescence. This minimum concentration decreases with the applied shear rate and with increasing molecular weight of the copolymer. Moreover, with the total molecular weight being the same, a copolymer with a larger PS block suppresses HDPE particle coalescence more efficiently than one with a smaller block. These results indicate that the interactions due to aforementioned steric repulsion of the presence of block copolymer at the interfaces are more important than those due to Marangoni stress.

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Intercalated, Exfoliated, and Dispersed Mechanism of Organoclay During Melt-Mixing

We now come to the mechanisms for the intercalation, exfoliation, and dispersion process of organoclay in polymer matrix during melt-mixing. Kawasumi et al. [79] prepared PP/organoclay nanocomposites in a TSE using the PP-o-g-MA as a compatibilizer to understand the miscibility effect of the PP-o-g-MA on the dispersibility of the layered silicate in the PP matrix. On the base of the investigation, they suggested a possible dispersion mechanism for polymer intercalation into the clay gallery spacings, as shown in Figure 6.9, which presents schematically the mixing process of the three components, i.e., PP, PP-o-g-MA, and organophilic

Stearyl ammonium

Silicate Layer of Clay

PP-o-g-MA Maleic Anhydride Group

PP

Figure 6.9 Schematic of dispersion process of organoclay in PP matrix with the aid of PP-o-g-MA. Reprinted from [79]. Copyright (1997) with permission from American Chemical Society.

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PSoz

Inatercalation Organophilic Clay Oxazoline Group

delamination

Figure 6.10 Schematic of organoclay dispersion process in PS/organoclay nanocomposites. Reprinted from [197]. Copyright (1999) with permission from John Wiley & Sons.

clay, into the nanocomposites. The PP-o-g-MA is first intercalated into the organophilic clay during melt blending to increase the interlayer spacing of the clay while the interaction of the layers should be weakened. The driving force for the intercalation originates from the maleic anhydride group and the oxygen groups of the silicates through strong hydrogen bonding. Then the intercalated structures are disordered by shear force during the shearing and the silicate layers are finally exfoliated into the PP matrix if the miscibility of PP-o-g-MA with PP is good enough to disperse the clay at the molecular level. A similar dispersion mechanism was proposed by Hasegawa et al. [197] for PS/organoclay nanocomposites prepared by simple melt-mixing of PSoz/PS and organophilic clays, where PSoz is a compatibilizer. Figure 6.10 is a schematic representation of the mixing process of PSoz and clay into the nanocomposites. It was suggested that the driving force of the hybridization originates from strong hydrogen bonding between the oxazoline groups of PSoz and the oxygen groups of the silicates. Based mostly on the experimental results, Dennis et al. [55] proposed a mechanism for dispersion of organoclay into a polymer during melt blending. The mechanism (Figure 6.11), based on the relationship between the compatibility of the chemistry of the clay treatment/the polymer matrix and the processing conditions, is described in three cases. The first case is chemistry-dependent. When the clay chemical treatment and the polymer are compatible, almost any set of processing conditions can be used to form an exfoliated nanocomposite. In case 2, clay chemical treatment and polymer are marginally compatible. In this situation, the processing conditions can be optimized to give an exfoliated nanocomposite. That is, the organoclay chemical treatment and the polymer are compatible enough that processing conditions can be tailored to optimize the exfoliation and dispersion. Finally, in case 3, there is no apparent compatibility of the clay chemical treatment and the polymer. Processing conditions can be optimized to give intercalants or tactoids that are minimized in size, but even partial exfoliation does not occur. Further, Fornes et al. [58] and Dennis et al. [55] proposed a model for organoclay exfoliation that envisions the role of both shear stress and residence time during melt-mixing of nanocomposites, as depicted in Figure 6.12. Initially, the stress should help break up large organoclay particles into dispersed stacks of

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8 μm Particle

Chemistry

Chemistry/Processing

Processing

Tactoids/ Intercalants

Dispersion

Partial Dispersion

Dispersion

Tactoids/ Intercalants

Tactoids/ Intercalants

Tactoids/ Intercalants

Figure 6.11 Schematic of mechanism for dispersion of organoclay into a polymer during melt blending. Reprinted from [55]. Copyright (2001) with permission from Elsevier.

Shear Stacks of silicate platelets or tactoids Organoclay particle (~8 µm)

(a)

Shear . Stress = η γ Shearing of platelet stacks leads to smaller tactoids (b) Shear

Diffusion Platelets peel apart by combined diffusion/shear process (c)

Figure 6.12 Stepwise mechanism of clay platelets exfoliation during melt mixing of nanocomposites: (a) organoclay particle breakup, (b) (intercalated) clay tactoid breakup, and (c) platelet exfoliation. Reprinted from [55, 58]. Copyright (2001) with permission from Elsevier.

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silicate tactoids, as shown in Figure 6.12a. Further down the extruder, transfer of the stress from the molten polymer to the silicate tactoids slides the platelets apart from each other and so decreases the height of platelet stacks, as illustrated in Figure 6.12b. Ultimately, as more polymer chains enter and go further in between clay platelets, especially near the edge of the clay galleries, individual platelets of the stacks peel apart from the edge because they are quite flexible and able to bend away from others in the stack, as suggested in Figure 6.12c. This does not require high shear intensity, but involves diffusion of polymer into the clay galleries driven by either physical or chemical affinity of the polymer for the organoclay surface and is thus facilitated by residence time in the mixing machine. This layer-by-layer peeling mechanism appears to greatly reduce the conventionally envisioned resistance to intercalation resulting from polymer chain confinement in the galleries between the platelets that are not able to bend as would be true in the mid-layers of a tactoid stack [77]. According to this model, an optimum balance between shear level and residence time is required to facilitate the exfoliation and dispersion of layered silicates. The shorter the stacks can be made by stress, the less time needed for peeling the layers away. For a fixed residence time, a higher stress should then translate into a higher degree of exfoliation. There may be, however, some critical stress needed to reduce the stack size. Moreover, peeling away the platelets of the stacks one by one requires the matrix polymer to have sufficient affinity for the clay surface to cause spontaneous wetting. This can be controlled by the nature of the organic treatment given to the clay to some extent. Stress alone cannot achieve exfoliation when the matrix polymer lacks sufficient compatibility with the organoclay. In the aforementioned melt-mixing process to prepare the nanocomposites, polymer chains first intercalate into the galleries of stacked layered silicates and the silicate layers then exfoliate into the polymer matrix. Hasegawa et al. [198] proposed a different mixing process for preparing the PA6/clay nanocomposites, in which a Na+ –clay (pristine clay) water slurry is used as an alternative for organically modified clay. The dispersion of exfoliated silicate layers of the Na+ –clay slurry into the PA6 matrix during mixing by a TSE occurs as follows (Figure 6.13). The clay slurry is first pumped into the melting matrix under vigorous shear

(a)

(b)

clay slurry (c)

(d)

nylon 6

silicate layer

Figure 6.13 Schematic of dispersion of silicate layers of Na+ –clay slurry into PA6 during melt-mixing. Reprinted from [198]. Copyright (2003) with permission from Elsevier.

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MMT swelling

H2O Injection

Increase of the interlayer distance

p-MMT

+ extr us io n

PA6 + H2O at 240°C, 100 bar Miscible system

Desorption of the water molecules

Diffusion and adsorption of PA6 Elimination of H2O

Shear

Exfoliated PA6/p-MMT nanocomposites

Figure 6.14 Schematic of dispersion of Na+ –clay in a PA 6 matrix during extrusion with water injection system. Reprinted from [199]. Copyright (2006) with permission from John Wiley & Sons.

(Figure 6.13a); the clay slurry reduces to finer drops during mixing and, at the same time, the water of the slurry drops begins to evaporate in contact with the matrix melt (Figure 6.13b and c); the evaporated water is removed under vacuum, and some silicate layers are fixed into melting matrix and disperse at monolayer, and some silicate layers gather and disperse at a few layers gathering (Figure 6.13d). By blending with the matrix vigorously and removing the water quickly, the exfoliated Na+ –clay silicate layers are fixed into the matrix almost as they are in water. In this process, the exfoliated silicate layers in water are directly fixed in the polymer matrix without aggregation of the silicate layers, which is different from the aforementioned mixing process using the organophilic clay. The results demonstrated that Na+ –clay silicate layers are exfoliated and dispersed homogeneously at nanometer level in the nanocomposites prepared using this novel process. Using a TSE with a water injection system, Fedullo et al. [199] prepared PA6/pristine clay (Na+ –clay or p-MMT) nanocomposites exhibiting exfoliated structure similar to nanocomposites prepared with organomodified clay. On this basis, they presented a model describing the mechanism of exfoliation/intercalation of the Na+ –clay in the PA6, as shown in Figure 6.14. During mixing, the injection of water enhances the fluidty and polarity of the PA chains and increases the intergallery distances of the Na+ –clay. The PA chains are able to diffuse between the platelets and are adsorbed on the silicate surface. This coating of the clay contributes to enhance the clay compatibility with the matrix, to reduce the electrostatic interaction between the platelets and as a matter of fact constitutes an in situ organomodification of the Na+ –clay. The problem of the dispersion and exfoliation of this Na+ –clay coated with PA chains is thus comparable to the dispersion of organomodified clay in a PA6 matrix.

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6.3 Material-Relevant Factors Affecting the Morphology The material-relevant factors that would affect the morphologies in multiphase polymer systems mainly include rheological and interfacial properties of the individual components, compatibilization, and composition. The rheological and interfacial properties of the individual components are critical factors for morphologies, and therefore final mechanical properties of multiphase polymer systems. Lee and Han [200] found that the viscosity ratio determines the state of dispersion for the symmetric blend composition and the blend ratio determines the state of dispersion for asymmetric blend compositions. Moreover, the size of dispersed phase in blends can generally be decreased obviously upon addition of nanoparticles, especially organoclays. 6.3.1 6.3.1.1

Viscosity of Components Binary Blends

The viscosity ratio (ηr ), which is defined as the ratio of the viscosity of the dispersed phase to that of the matrix phase, has been shown to be one of the most critical factors for controlling the binary blend morphology [169]. In binary blends, when both the dispersed and matrix phases are viscoelastic, it is generally believed that the ηr should be close to unity in order to obtain the finest dispersion. Relatively small and uniformly dispersed phases are usually advantageous for the properties of polymer blends. For example, Cheng and Wang [201] adjusted the ηr to control the morphology of TPU in the POM matrix by changing the processing temperature. The results showed that a highly oriented filamental morphology developed at a ηr of about unity is most advantageous for the improvement of the toughness of the blends. Highly viscous matrices (ηr  1) can enhance the droplet breakup owing to their efficient shear stress transfer towards the dispersed phase and to a somewhat reduced coalescence probability during melt-mixing; whereas low viscous matrices (ηr  1) often act as a lubricant for the dispersed phase hampering the droplet breakup due to insufficient stress transfer to the dispersed phase [181, 202]. With regard to the flow of binary blends, it is generally acknowledged and observed that a lower-viscosity component tends to encapsulate a high-viscosity component, with the effect of the lower-viscosity component forming the continuous phase over wider component composition ranges or at a much lower volume fraction than the higher-viscosity component [203, 204]. Grace [134] determined experimentally the Cacrit versus the ηr for a single Newtonian drop suspended in a Newtonian medium in both shear and elongational flows. The Cacrit is as follows in shear flow: log Cacrit = −0.370 + 0.185(log ηr )2 + 0.0177(log ηr )3 −

0.115 log ηr − log 4.5

(6.9)

The results by Grace [134] showed that drops deform into ellipsoids under a shear flow, but do not break at a ηr > 3.7. A variety of studies have shown, however, that the ηr requirement for processing polymer blends can be much broader than expected. After studying several PA66/EPR and PET/EPR blends with 15 wt% EPR phase, prepared using a TSE, Wu [3] concluded that the formation of dispersed phase can be described by a master curve between the Ca and the ηr and proposed the following equation: Dn =

4σ ±0.84 η ηm γ˙ r

(6.10)

where Dn is the predicted number-average diameter of dispersed phase. The plus (+) sign in the exponent applies for ηr > 1 and the minus (–) sign for ηr < 1. Wu [3] showed that the dispersed EPR phase in

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PA66/EPR blends can break up even when ηr > 4. It should be noted that Wu’s equation takes into account the coalescence process, but not the composition effect. Moreover, this is an empirical correlation and has no theoretical basis. It provides, however, a good first estimation of droplet size for polymer blends. Eq. (6.10) indicates that the droplets are smaller when the interfacial tension is lower, the shear rate is higher, and the ηr is closer to unity. Several studies illustrated that a linear relationship exists between the dispersed phase size and the ηr . Favis and Chalifoux [205] found that the ηr has a significant effect on the phase morphology of PP/PC blends with the phase size increasing by a factor of 3 to 4 times when increasing the ηr from 4.5 to 17.3. An increase in the dispersed phase size is observed at ηr below 1 for blends with the PP as the matrix. A minimum dispersed phase size is obtained at a ηr of about 0.15 in blends with the PP as the dispersed phase. Hietaoja et al. [206] also found a linear increase in the particle size with increasing ηr for PP/PA66 blends with different PPs as matrixes, but with no correlation between them when using PP as dispersed phase. The results from previous investigations, however, do not always agree due to differences in the content of the dispersed phases, mixing technologies, particle size measurements methods, and shear rates selected for the calculation of the ηr . For example, P¨otschke et al. [202] found that the increase in the particle size with the ηr for TPU/PO blends is higher than that reported for PP/PC [205] but lower than that found for PP/PA6 [3] blends. This results from the different interfacial properties and blending technologies used. Some works have been carried out to investigate the morphology of blends with low [204] or extremely low [93, 207] and high [179] or extremely high [88, 208, 209] ηr . Using PC/PE, amorphous copolyester (PETG)/PE, and PB/PCL blends, the studies by Ratnagiri and Scott [204] showed that when the ηr is less than about 0.2, phase inversion during mixing occurs in blends with the dispersed phase at only 10 wt%, independent of the relative transition temperatures. Leblanc et al. [93, 207] prepared immiscible PA66/PEG blends exhibiting ηr in the magnitude of 10−5 and found that the PEG concentration, shear rate and ηr clearly appear to be the key parameters affecting the morphology. Work performed by Huneault et al. [208] showed that PP/HDPE blends with a ηr up to 1000 can be dispersed in the micron range and that the real challenge in that case is achieving blending homogeneity and reducing the concentration of gels. Chen et al. [88] investigated the phase morphology development in the PEG/PA12 blend with an extremely high ηr (in the magnitude of 102 –103 ). The results showed that the breakup process of the dispersed PA12 phase is observed for this blend. The PA12 forms the spherical particles in the PEG matrix. Many small droplets with a diameter of 0.1–10 μm develop in the initial stage of mixing, which demonstrates that the major breakup of pellets occurs at the very beginning. The volume average size of the dispersed phase increases with the time in the initial period and then levels off. Its size distribution narrows with prolongation of the mixing time and tends to be a single narrow peak. Rameshwaram et al. [209] added an ultrahigh viscosity polymer (PB) to the matrix (PDMS) with a very low viscosity. This results in the formation of polymer nanocomposite-like blends with an ultrahigh ηr (162,000) and its viscoelastic behavior is similar to those of polymer/clay nanocomposites. This is probably because the continuous PDMS phase tends to break up into small droplets and diffuse into the large PB domains. Very recently, Mukherjee and Sarkar [210] numerically simulated the effect of ηr on a viscoelastic drop deforming in a Newtonian matrix under steady shear. It was found that for the ηr lower than unity, the longtime steady deformation behavior is similar to that of the viscosity matched system: the drop shows reduced deformation with increasing Deborah number (De) due to the increased inhibiting viscoelastic normal stress inside the drop. However, for higher ηr systems, the drop deformation displays a non-monotonic behavior with increasing De: the deformation first decreases, but above a critical De, it increases. Some different viewpoints have been suggested. Kang et al. [211] studied the nature of initial mixing of a two immiscible fluid system in a two-dimensional rectangular cavity to understand the drop breakup mechanisms induced by shear in a fully transient flow field typical of a SSE. The results showed that the transient breakup process is more dependent on the matrix viscosity than on the ηr , with a higher matrix

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viscosity resulting in a more evenly distributed dispersion phenomenon by delaying the breakup time to inhibit the premature relaxation in the downward flight region. Most reported studies used the ηr data generated at the same shear rate. Namhata et al. [203] developed a model to predict the blend morphology development during melt-flow based on the fact that when two phases with different viscosities are subjected to a stress field, each phase has a different velocity resulting in a difference in volumetric flow rate. They stated that the ηr should be determined at the same shear stress instead of using a ηr criterion based on viscosities measured at the same shear rate, especially at higher shear conditions. 6.3.1.2

Ternary Blends

In ternary polymer blends in which three phases coexist, the definition of viscosity ratio is a critical topic [212]. Reignier et al. [213] showed that in order to study the influence of viscosity ratio on the morphology of multiphase blends, the viscosity ratio should be estimated at a constant shear stress rather than at a constant shear rate, because the shear stress, rather than the shear rate, is continuous at the interface between the dispersed phase and the matrix phase. Kim et al. [214] reported that for the PP/rubber (EPM or EPDM)/PE ternary blends, when the two dispersed phases have equal compositions, the dispersed phase with lower viscosity forms a shell encapsulating the one with higher viscosity. Gupta and Srinivasan [35] observed that in PP/SEBS/PC blends, when SEBS has a higher viscosity than PC, SEBS forms a boundary layer around the PC phase, and otherwise separated disperse morphology is observed. Based on the deviation of viscosity–composition curves from the log-additivity rule, Utracki et al. [215] divided the curves into three types: positive-deviation blends (PDB), negative-deviation blends (NDB), and positive-negative-deviation blends (PNDB). Ha et al. [216] plotted complex viscosities versus blend composition for PP/mPE/HDPE ternary blends (Figure 6.15). The curves showed a minimum at 25 wt%

0.1 rad/s 1.0 rad/s

η* (poise)

10 rad/s 104

100 rad/s

398.1 rad/s

103

0

20

40 60 mPE/(HDPE+mPE) (wt%)

80

100

Figure 6.15 Plots of complex viscosity (η*) versus HDPE/mPE binary composition in PP/mPE/HDPE ternary blends at 230 ◦ C. Reprinted from [216]. Copyright (2004) with permission from John Wiley & Sons.

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exhibiting NDB type at all frequencies. This is attributed to the fact that the viscosity of the PP matrix is much lower than that of the composite droplet dispersed phase of HDPE/mPE blends, as well as HDPE. Ha et al. [216] believed that the phase inversion occurs at this inflection point. The effect of the viscosity ratio on the phase morphology of ternary blends is very complex, especially in the case of a core–shell structure for the dispersed phase [217]. Do et al. [218] prepared iPP/EPR/HDPE (80/10/10 wt) ternary blends using a co-rotating TSE. Rubber domain (EPR+HDPE) size can be easily controlled by the addition of HDPE. When HDPE satisfying ηHDPE < ηEPR is added, the rubber domain size decreases, and vice versa. The study by Abolhasani et al. [219] demonstrated that a core–shell type morphology is developed for immiscible ternary blends of PET/EVA/PP (PET as the matrix and (PP/EVA) composition ratio = 1/1). When the viscosity and elasticity of the shell (EVA) are less than those of the core (PP), the shell phase dominantly controls the droplet behavior. However, when the viscosity and elasticity of the shell are higher than those of the core, the mentioned phenomena is not observed. Luzinov et al. [34] described a core–shell morphology for PS/SBR/PE blends with different viscosity ratios of components. They demonstrated that the core diameter is influenced by the viscosity ratio of core phase to shell phase whereas the composite droplet size is affected by the viscosity ratio of the shell phase to matrix. Hemmati et al. [217] modified this concept and used the ratio of average viscosity of two minor phases to the viscosity of matrix to predict the dispersed phase size. Average viscosity (ηav ) is calculated by a simple mixture rule: ηav = η1 φ1 + η2 φ2

(6.11)

where η1 and η2 are viscosities of minor phases, φ 1 and φ 2 are their volume fractions. Ha and Kim [220] showed that in PP/mPE/HDPE blends, fibrils are observed when the viscosity of the dispersed-phase (mPE/HDPE) is lower than that of the PP matrix, or when the viscosity of mPE is lower than that of PP, although the dispersed-phase viscosity is higher than the PP viscosity, indicating that the mPE/HDPE blends have a core–shell morphology. With a core–shell morphology, some mPE in the shell forms fibrils. For the same ternary blends, Ha et al. [216] found that when the mPE encapsulates the HDPE in the PP matrix, compared to the encapsulation of mPE by HDPE, better mechanical properties are obtained, presumably because the compatibility between PP and mPE is better than that of PP and HDPE. The results regarding the effect of viscosity ratio of the two minor phases in a ternary blend on the resultant morphology are, however, controversial. Some studies suggested that composite droplet formation may be related to the viscosity of the dispersed components. Gupta and Srinivasan [35] showed that for PP/SEBS/PC blends with the PP as the matrix, the SEBS component forms a boundary layer at the surface of the PC droplets when PC is less viscous than SEBS, whereas the minor components are separately dispersed when PC is more viscous than SEBS. Nemirovski et al. [221] observed the reverse effect for a number of ternary thermoplastic/thermotropic blends. They suggested that dispersed component A encapsulates dispersed component B in matrix C when both the thermodynamic effects, expressed by means of a positive spreading coefficient λAB , and kinetic effects, expressed by means of dispersed phase viscosity ratio (ηA /ηB ) smaller than 1, act cooperatively. However, in some cases, a value of ηA /ηB greater than 1 hinders the development of the core–shell structure (A encapsulating B) even though encapsulation is predicted by a positive λAB . However, most researchers did not find any influence of the dispersed phase core–shell viscosity ratio on the composite droplet structure [217, 221]. Luzinov et al. [222] reported that for ternary blends consisting of PS, SBR and different POs, changing the SBR/PO dispersed phase viscosity ratio from 0.4 to 4.5 has no effect on the development of the core–shell morphology, the PO phase being systematically encapsulated by SBR as predicted by the spreading coefficient analysis. Hemmati et al. [217] showed for various 70/15/15 PP/rubber/HDPE ternary blends that changing the rubber/HDPE viscosity ratio from 0.86 to 100 has no effect on the morphology; in all cases the HDPE component is encapsulated by the rubber phase, in agreement with the theoretical prediction based on the spreading coefficient.

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Nanocomposites

Some studies regarding the effect of the viscosity or molecular weight (MW) of polymer matrix on the microstructure of nanocomposites prepared by melt-mixing have produced contradictory results. For PS/clay nanocomposites prepared by melt intercalation, Vaia and Giannelis [78] found that the MW of PS only appears to affect the kinetics of polymer intercalation but does not affect the final hybrid structure. The results by Fornes et al. [58] showed that the MW of PA6 matrix has a significant effect on the microstructure and mechanical properties of nanocomposites, which were prepared based on three different MW grades of PA6 (low, medium, and high) using a co-rotating TSE. WAXD and TEM results collectively reveal well exfoliated structures for the MMW and HMW PA6s based nanocomposites, regardless of the organoclay used; whereas the LMW PA6 based nanocomposites exhibit a mixed structure, having regions of intercalated and exfoliated clay platelets. The average number of platelets per stack is shown to decrease with increasing PA6 molecular weight, thereby revealing larger extents of clay platelet exfoliation for the nanocomposites in the order HMW > MMW > LMW nanocomposites. It is believed that the role of polymer molecular weight is attributed to the shear stress. The melt viscosity of the three systems follows the order HMW > MMW > LMW, and hence the resulting shear stresses exerted on the organoclay by the polymers also follow the same order. Therefore, the levels of stress by higher MW PA6s are significantly higher during melt-mixing. As displayed in Figure 6.12, higher levels of shear stress can shear the taller stacks into shorter ones and aid in the breakup of clay particles, ultimately improving the clay platelet exfoliation, whereas the LMW PA6 melt may not generate stresses above a critical level needed to reduce stack size. Moreover, the LMW based nanocomposites have larger clay platelet lengths than the higher molecular weight ones. There are several possible reasons for the relationship between polymer molecular weight and platelet length. The shorter lengths for the higher molecular weight based nanocomposites may be a result of attrition due to higher melt viscosities of the PA matrix. Melt viscosity differences among the LMW, MMW, and HMW matrices may also play another role in determining the observed platelet length. The higher molecular weight matrices, due to their higher melt viscosities, transfer more stress or energy to achieve separation of platelets. However, a low molecular matrix imparts low shear stresses on the agglomerates, which may skew the stack of platelets rather than separate them, as displayed in Figure 6.16. Such skewed stacks would appear larger. 6.3.2

Elasticity of Components

The fundamental understanding of droplet shape and breakup is based on Taylor’s theory [115, 133]. However, Taylor’s theory (Eqs. (6.1) and (6.2)) is hardly applicable to viscoelastic systems, such as polymer blends, for which the elasticity (particularly the flow-induced elasticity ratio) of the blend components has a strong effect on the dynamic equilibrium between droplet breakup and coalescence [102, 104, 108, 223] and so on the final

Figure 6.16 A possible mechanism for the cause of larger effective particle sizes for LMW PA6 based nanocomposites. Reprinted from [58]. Copyright (2001) with permission from Elseiver.

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phase morphology [4, 27, 101, 108, 179]. That is, the mechanism of droplet deformation and breakup is quite different in viscoelastic systems from that in Newtonian systems. When one or both of the components are viscoelastic in a binary blend, the fluidodynamics of the droplet becomes more complex, as the constitutive time scales of the two fluids also come into play, together with the intrinsic time scale related to the very existence of an interface [114]. Apart from its correlation with the ηr , the Cacrit for viscoelastic systems is additionally influenced by component elasticity [17]. To understand the effect of the component elasticity on the droplet deformation and breakup, extensive theoretical and experimental works have been performed. A classical study about the role of the component elasticity in influencing the blend phase morphology was performed by VanOene [2] for capillary flow of PS/PMMA blends. It was demonstrated that under the conditions of dynamic flow, the elasticity differences between the components of a blend can alter the interfacial tension, which can be quite different from that in the absence of flow. He proposed an expression as follows: σd = σs +

D (N2,d − N2,m ) 12

(6.12)

where σ d and σ s are the dynamic and static interfacial tensions, respectively; N 2 is the second normal stress difference; Subscripts d and m represent the dispersed and matrix phases, respectively. Reignier et al. [213] believed that the N 2 in Eq. (6.12) is equivalent to the current definition of N 1 (first normal stress difference) and so rewrote VanOene’s equation as follows: σd = σs +

D (N1,d − N1,m ) 12

(6.13)

In this form, one can see that the VanOene’s equation qualitatively shows that when (N1,d − N1,m ) is positive, the elasticity acts as an additional interfacial tension under shear conditions; whereas in the opposite case, the interfacial tension is reduced under shear. When considering the effect of component elasticity, the deformation and final morphology of the dispersed phase in polymer blends are the result of a dynamic equilibrium between the deforming forces and the deformation-resisting forces. The deforming forces are from the matrix, which consist of the shear stress and the matrix first normal stress (T 11,m ). The deformation-resisting forces include the interfacial tension and the droplet first normal stress (T 11,d ). Since T 11 is much greater than T 22 (second normal stress) [101], N 1 (= T 11 –T 22 ) can be approximated by T 11 [104]. Then the following equation can be obtained: ηm γ˙ + N1,m ∝ 2σ/D + N1,d

(6.14)

On this basis, Seo and Kim [224] defined a new capillary number–elastic capillary number, CaE , as: Ca E =

N1,m − N1,d 2σ/D

(6.15)

That is, CaE is the ratio of the elastic force on the droplet to the interfacial tension. Dispersed particles are deformed when CaE > 1. In addition, based on the simple force proportionality mentioned above, Ghodgaonkar and Sundararaj [104] predicted the diameter of a single polymer droplet in polymer blends considering the contribution of the viscoelastic properties of polymers. At low shear rates, N 1 can be approximated by 2G  , then Eq. (6.14)

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can be rewritten as follows: ηm γ˙ + 2G m ∝ 2σ/D + 2G d

(6.16)

At higher shear rates (>10 s−1 ), N 1 > 2G  , but the two quantities are proportional to each other. As a first approximation to obtain qualitative behavior, the two sides of Eq. (6.16) are equated to obtain an equation for droplet diameter: D=

2σ ηm γ˙ − 2(G d − G m )

(6.17)

Eq. (6.17) cannot be used when the T 11,d is significantly greater than the T 11,m . For the case where only the droplet elasticity is considered, Eq. (6.17) can be simplied as follows: D=

2σ ηm γ˙ − 2G d

(6.18)

The predicted diameters of the dispersed phase from Eqs. (6.17) and (6.18) are comparable to the experimentally determined diameters for several blends at different operating conditions [104]. It was observed for some blends (PS/PP, PS/EPMA, PS/amorphous PA) that, as the shear rate increases, the dispersed phase diameter initially decreases. At a critical shear rate, the diameter reaches a minimum value, and beyond it, the diameter increases with shear. This critical value was found to be between 100 to 162.5 s−1 for a PS/PP blend (ηr = 0.87) with PP of less than 8 wt%. This minimum was not observed for the blends with higher dispersed phase fraction, such as 80/20 PS/PP blend. This may be due to greater coalescence. Some experimental studies were carried out for blends in which only one component is viscoelastic. Gauthier et al. [225] studied the deformation and breakup of Newtonian droplets suspended in a viscoelastic liquid and of viscoelastic droplets in a purely viscous Newtonian liquid in Poiseuille flow. They found that at low values of ηr , viscoelastic droplets behave in a manner similar to Newtonian ones. However, at higher values of ηr , the droplets are pulled into threads as they breakup, and the resulting Cacrit is much larger. Flumerfelt [1], who studied the breakup of isolated Newtonian droplets in viscoelastic polymer fluids in a simple shear flow field, found that there exists a minimum droplet size below which breakup cannot be achieved. This minimum droplet size is larger when the matrix is elastic. Elmendorp and Maalcke [135], investigating the contribution of elasticity on the breakup of isolated viscoelastic droplets in Newtonian matrices and of Newtonian droplets in viscoelastic matrices in a simple shear flow, found that more elastic droplets (as measured by the N 1 ) are more stable against breakup, whereas the more elastic matrices lead to increasingly unstable droplets. Varanasi et al. [110] utilized a viscoelastic Boger fluid for the isolated droplet and a purely viscous Newtonian fluid for the matrix and so could separate the effects of the ηr and elasticity. They observed a linear relationship between the Cacrit and the calculated N 1 of the droplet at a fixed ηr . Most polymer blending involves viscoelastic dispersed phases in a viscoelastic matrix. So, blends have also been investigated in which both the dispersed and matrix phases are viscoelastic. Using PA66 or PS as the matrix phase and EPR as the dispersed phase (15 wt%), Wu [3] found that the steady-state Ca or steady-state droplet size for these extruded viscoelastic blends is generally about 10 times larger than that for Newtonian blends at the same ηr and shear rate. Mighri and co-workers investigated the influence of elasticity ratio, as defined as the Maxwell relaxation time ratio between the droplet and matrix, λd /λm , on isolated droplet deformation in an uniaxial elongational flow [5] and on isolated droplet deformation and breakup in a shear flow [6]. Here, λ = N1 /2ηγ˙ 2 . In these studies, both the droplet and matrix consist of constant-viscosity elastic (Boger) fluids. They found that the Cacrit for viscoelastic droplets increases with increasing droplet

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elasticity, but decreases with increasing matrix viscoelasticity. In shear flow, the Cacrit increases drastically with increasing λd /λm up to λd /λm equal to 4, attains a maximum of about 1.75 at λd /λm of 4, but does not change much further for λd /λm > 4. The droplet deformation diminishes as λd /λm increases, in either the elongational or shear flow, which is consistent with the findings of Elmendorp and Maalcke [135]. They further noted that for high matrix elasticity (λd /λm ≤ 0.37), the deformation of elastic droplets in an elastic matrix is larger than that of Newtonian droplets in a Newtonian matrix with the same ηr and interfacial tension. However, for λd /λm > 0.37, the elastic droplets deform less than a Newtonian droplet in a Newtonian matrix. In elongational flow, the droplet deformation is more affected by matrix elasticity for λd /λm ≤ 0.2, whereas by droplet elasticity for λd /λm > 0.2. Using 80/20 HDPE/PS blends with a ηr of 0.5, 1 and 2, Lerdwijitjarud et al. [9] investigated the influence of elasticity ratio on steady-state dispersed-phase size. The elasticity ratio was measured by the droplet/matrix first normal stress difference ratio, N 1r = N 1,d /N 1,m . The viscoelastic droplets in a viscoelastic matrix are more difficult to break compared to Newtonian droplets in a Newtonian matrix, because of the contribution of both droplet elasticity and matrix shear-thinning. For all blends studied, the steady-state capillary numbers (i.e. dimensionless droplet sizes) increase with N 1r and follow a power law with scaling exponents between 1.7 and 1.9. These capillary numbers are 4–80 times higher than the Cacrit for the system of a Newtonian droplet in a Newtonian matrix. Lerdwijitjarud et al. [106] studied the deformation and breakup of droplets of weakly elastic PBd Boger fluid (Wid ≤ 0.02) sheared in a Newtonian matrix at fixed ηr . Here, Wi is Weissenberg number, a ratio of elastic to viscous forces, which can be estimated using either the N 1 or G  as a measure of the elastic forces and the shear stress as a measure of the viscous forces. Both isolated PBd droplets and 20 wt% PBd droplets in a Newtonian PDMS matrix were considered. They found that droplet elasticity results in a reduction in the degree of droplet deformation at any given shear rate and a corresponding slight (up to about 20%) increase in the Cacrit compared with a Newtonian droplet. However, the steady-state droplet size in the 20 wt% dispersed phase blend is smaller than that in isolated droplets under identical shearing conditions. This result suggests that local increase in shear stress present in concentrated blend is more important than coalescence in influencing steady-state droplet size. Later, Lerdwijitjarud et al. [13] investigated the same systems with a fixed ηr of about unity as those in previous works [106] but with a much higher droplet elasticity (with Wid ranging up to about 3). A linear relationship between the Cacrit and the Wid holds up to a value of Wid about unity, with a saturation of Cacrit at about 0.95 for high Wid . Using in situ microscopic visualization, Mechbal and Bousmina [226] studied the droplet deformation on model systems of PIB/PDMS during mild shear flow. The obtained results showed that elasticity of the droplet decreases its deformation and increases the wobbling, whereas the elasticity of the matrix increases the droplet deformation without inducing wobbling. Most studies, including the aforementioned works and some other works [4, 101, 104, 124, 161], show that the matrix elasticity enhances the droplet deformation, causing the droplet to break at a lower Ca; whereas the droplet elasticity hinders the droplet deformation, making it break at a higher Ca. These trends are supported by theoretical considerations which suggest that the droplet elasticity should augment the interfacial tension making droplets less deformable and less breakable. The matrix elasticity, on the other hand, should produce high extensional stress near the stagnation point at the downstream end of the droplet leading to tension that tends to stretch the droplet, making it easier to break [9]. Moreover, viscoelastic effects within the droplet inhibit its deformation not only by their contribution to the viscous and elastic stresses but also by affecting the flow field [117, 161]. However, there have been conflicting theoretical and experimental reports about the effect of component elasticity on droplet deformation or phase morphology. Milliken and Leal [177] investigated the deformation and breakup of isolated viscoelastic droplets in a Newtonian matrix fluid under two-dimensional elongational flows. When compared to Newtonian droplets with equivalent viscosity ratios, they observed greater droplet deformation for non-Newtonian droplets at a given Ca. They also observed two modes of droplet breakup, tip streaming and tip streaming with stretch, that differ substantially from the

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observed breakup of Newtonian droplets. Ultracki and Shi [105] reported that droplet elasticity reduces the droplet deformation for the ηr below 0.5, but increases the deformation for the ηr above 0.5. Varanasi et al. [110] showed that for any fixed ηr there exists a critical shear rate, below which the viscoelastic droplets sheared in viscous Newtonian fluid are easier to breakup than purely viscous Newtonian droplets. Sibillo et al. [114] explored the effect of matrix elasticity on the breakup of an isolated Newtonian droplet under step shear flow. Three viscosity ratios were considered, i.e. 2, 0.6 and 0.04. They found that at all these viscosity ratios, droplet breakup is hindered by the matrix elasticity. The results of Flummerfelt [1] and Guido et al. [113] also lend support to the findings of Sibillo et al. [114]. Verhulst et al. [227], studying the effect of matrix elasticity on the steady and transient deformation of a Newtonian droplet in shear flow, demonstrated that matrix elasticity reduces the steady droplet deformation and promotes droplet orientation, can induce a droplet deformation overshoot in the start-up experiments and slows down the relaxation phenomena. The component elasticity may also contribute to phenomena such as sheet formation [101], widening of droplets [4] or elongation [27, 108, 179] in the vorticity direction (as shown in Figure 6.6). For example, remarkable effects of component elasticity have been observed along the vorticity direction. A variety of systems, flow types, and experimental conditions give rise to this phenomenon [4, 8, 17, 107]. It is generally accepted that alignment and breakup of droplets along the vorticity direction involve a complicated interplay among the first and second normal stress differences of both phases, shear-thinning behavior, ηr , and applied Ca [31]. Levitt et al. [4] traced PP droplets with varying viscosity and elasticity sheared in a high-elasticity PS matrix. For low ηr systems, they visualized a remarkable widening of PP droplets along the vorticity direction. The authors attributed this transient vorticity widening to the normal stresses exerted by the matrix on the less elastic droplets, which overcomes the contraction in the opposite direction caused by the interfacial tension. The width of the flattened droplets depends on the difference in elasticities between matrix and droplet and is proportional to the N 2 between two phases. Levitt et al. [4] assumed that the stretch in the hoop direction is much larger than that in the thickness direction and that the N 2 is proportional to the N 1 , then estimated N 2 from easily measured G  and finally approximated half the maximum thickness (bmax ) of the flattened droplets as follows: bmax =

σ σ ≈  0.6(N2,d − N2,m ) G m − G d

(6.19)

Levitt et al. [4] also showed that the sheet formation and the widening phenomena result in the area generation larger than the value calculated for the affine deformation. However, for a ηr larger than 1 and elasticity ratio above 2, no droplet widening is observed. The deformation of PS droplets inside an HDPE matrix was visualized using a pressure-driven optical flow cell situated at the exit of a TSE by Migler et al. [27] and at the die region by Hobbie and Migler [179]. Different dilute HDPE/PS blends with ηr up to 240 were studied. The droplets are transformed from ellipsoids at low Ca to strings at modest Ca and back again to spheres of the same size as found under low Ca. Droplet viscoelasticity is believed to play a role in this transition. They also observed a variety of droplet shapes as a function of shear rate and ηr . At a higher shear rate (about 280 s−1 ) and a ηr of 1.8, PS droplets are elongated along the vorticity direction, which is believed to be ascribed to the difference between droplet and matrix elasticity. By extrapolating data at high shear rates to lower rates, the authors obtained the Cacrit for droplet vorticity elongation of about 53, 13, and 11 for the ηr of 1.8, 22, and 240, respectively [179]. The increase in the Cacrit for this vorticity elongation with the ηr is consistent with a mechanism involving droplet elasticity, since a higher droplet viscosity would require a higher external shear rate to attain the same internal shear rate within the droplet, which would be needed to maintain a high elasticity of the droplet fluid. This vorticity alignment was also observed by Migler [108], who considered a case of highly elastic PIB droplets deformed

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in PDMS matrix under shear. In this system, the ηr is near unity, but the elasticity ratio of the dispersed phase to the matrix is greater than 100. He found that under conditions of weak shear and small droplets (Ca < 5), the droplets align along the shear direction, whereas for strong shear and large droplets (Ca > 5), the alignment is along the vorticity direction with a broad distribution of aspect ratios of droplets. Subsequently, Mighri and Huneault [8, 96] investigated the role played by the viscoelastic behavior of droplet/matrix phases on the dispersion mechanisms using model (Boger fluid in PDMS) viscoelastic systems and polymer (PE/PS and PP/EPR) systems. The deformation and breakup of these viscoelastic systems present two main differences with Newtonian systems. First, above a critical shear rate (Ca ∼ 5), the deformed droplet begins to contract in the flow direction and align in the vorticity direction. This vorticity alignment increases with further increasing the shear rate until the final breakup occurs at a Ca no higher than about 35. The authors proposed that the critical shear stress for this vorticity alignment is probably related to the values of the first and second normal stress differences and their dependencies on shear rate. This vorticity alignment dynamics is directly related to the elastic normal stress inside the droplet that results from a flow-induced circulatory flow in the droplet. Secondly, the breakup occurs when the ends of highly elongated particles get in slightly different planes. Using HDPE as matrix and PS as droplet, Cherdhirankorn et al. [107] investigated the dynamics of deformation of an elastic droplet in an elastic matrix by selecting two blends with ηr near unity, but of different elasticities of both droplet and matrix phases with the Wim of about 0.1–0.5 and Wid of about 0.2–0.5. The Wid was characterized by the droplet fluid properties at the macroscopic shear rate measured on the pure droplet fluid in a rheometer. In start up of a steady shear flow, the different elasticities in the polymer blends lead to qualitative differences in the droplet deformations that occur before the droplet reaches its steady-state shape. In the blend with higher elasticity, the deformation of the droplet oscillates several times before attaining its steady-state shape. In the blend with lower elasticity, the droplet first deforms in the shear direction, and thereafter continuously contacts in the flow direction until it reaches its steady-state shape. When the Ca is increased at fixed shear rate (and hence fixed elasticity) by increasing the droplet size in both blends, the steady-state droplet shape becomes increasingly elongated in the vorticity direction and develops cusps along the vorticity axis. In the blend with higher elasticity, at still higher Ca, droplet breakup occurs when two ends of a droplet elongated in the vorticity direction are situated on streamlines of different velocity which pull the droplet ends apart, leading to rupture. The aforementioned investigations focus on the effect of component elasticity on the deformability of the dispersed phase in binary immiscible blends. Reignier et al. [213] proposed a dynamic interfacial tension model (Eq. (6.23)) to investigate the influence of elasticity on composite droplet formation for HDPE/PS/PMMA ternary blend. The results showed that encapsulation effects in composite droplet type systems can be predicted based on this model. 6.3.3

Interfacial Tension

Most of the research on the subject of multiphase blends showed that their morphology can be predicted through the knowledge of interfacial tension between the components [212, 228]. Keeping all other parameters constant, increasing the interfacial tension is expected to have two effects. First, it increases the dispersed phase sizes. Second, it decreases the value of the capillary number, reducing the region of distributive mixing and thus reducing the range of co-continuity [229]. Some studies [187, 230] indicated unambiguously that a direct relationship exists between dispersed phase size and interfacial tension in the absence of coalescence effects for blends prepared using an internal mixer. These results demonstrate a direct experimental confirmation of the phase size/interfacial tension relationship predicted by Taylor theory (Eqs. (6.1) and (6.2)). The results by Huneault and co-workers [96, 208], however, seemingly indicated that the interfacial tension plays a much smaller role in the breakup for highly viscous polymer blends, such as PP/PE, PE/PA, and PP/EPR. Moreover,

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interfacial tension relaxes the morphology by three mechanisms: retracting nonspherical droplets toward a spherical shape, breaking up extended droplets into smaller spheres, and coalescing small droplets to make larger ones [10]. Under equilibrium mixing conditions, the phase morphology of a multiphase blend may be predicted from the balance of interfacial tensions between the components. By using the concept of a spreading coefficient, Hobbs et al. [37] rewrote the Harkin’s equation in which two distinct polymer phases are dispersed within a matrix by substituting the appropriate interfacial tensions for the surface tension values: λij = σjk − σik − σij

(6.20)

where σ jk , σ ik , and σ ij are the interfacial tensions between different component pairs; λij is defined as the spreading coefficient describing the tendency of component i to encapsulate or spread around component j in a matrix of component k to minimize the interfacial free energy of the overall system. In fact, Eq. (6.20) (is called as spreading coefficient model) conveys, for a ternary blend, the tendency for the dispersed phase which forms the highest interfacial tension with the matrix to enter within the other dispersed phase [94]. In ternary blends consisting of one major constituent (component A) and two minor ones (components B and C), four different morphologies are possible (Figure 6.17): (a) separate dispersions; (b) partial encapsulation; (c, d) two possibilities of complete encapsulation (composite droplets). Each morphology has a distinct set of spreading coefficients. λBC must be positive for component C to be encapsulated by component B. When both λBC and λCB are negative, components B and C tend to be dispersed separately within matrix A. Furthermore, in the particular case where λAC is also negative, a partial encapsulation morphology occurs. In the intermediate region, where λBC ≈ 0, a ‘stack’ [231] or ‘acorn’ morphology may result (Figure 6.18), in which component B only partially eliminates the interface between component C and matrix component A. Guo et al. [33] extended the spreading coefficient approach to take into account the overall surface free energy by including the interfacial area of each component. They developed the following equation which states that the equilibrium phase structure of a multiphase system is determined by the lowest free energy

λBC < 0

λBC < 0

λBC > 0

λBC < 0

λCB < 0

λCB < 0

λCB < 0

λCB > 0

λAB > 0 or λAC > 0

λAB < 0 or λAC < 0

λAB < 0

λAB < 0

(c)

(d)

(a)

(b)

Figure 6.17 Schematic of possible morphologies in a ternary blend composing of one major phase (phase A (white)) and two minor ones (phases B (gray) and C (black)): (a) separate dispersions; (b) partial encapsulation; (c) and (d) complete encapsulation (composite droplets). Reprinted from [228]. Copyright (2006) with permission from American Chemical Society.

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Figure 6.18 An ‘acorn’ dispersed phase morphology forming when λBC ≈ 0. Reprinted from [232]. Copyright (2004) with permission from Elsevier.

state defined by: G=



n i μi +



Ai σi j

(6.21)

i = j

i

where G is the Gibbs free energy, n is the number of moles, μ is the chemical potential, and A is the interfacial area. Eq. (6.21) can be called as minimum free interfacial energy model. The first term on the right of Eq. (6.21) is equal for all types of morphologies, but the second term differs from one morphology to another. The lowest value of the second term corresponds to the most stable morphology. This model predicts that the interfacial tensions play the major role in developing the phase morphology, whereas a less significant, but still important, role is played by the surface areas of the dispersed phases, as determined by the composition. For an N-component system, there are (N – 1) interfaces coexisting in the system at the lowest free energy. For a ternary system, there are two co-existing interfaces. Because the first term on the right of Eq. (6.21) is the same, it can be ignored. To further simplify the problem, the surface areas of the minor phases, AB and AC , can be calculated based on average phase sizes. For three phase morphologies as shown in Figure 6.17a, c, and d, that is, components B and C form separate phases (B + C), component C is encapsulated by component B (B/C), and component B is encapsulated by component C (C/B), the interfacial free energies of the system can be calculated by using the following equations:   

Ai σi j Ai σi j Ai σi j

  B+C  B/C C/B

1/3

1/3

= (4π )1/3 (n B x 2/3 σ AB + n C σ AC )(3VC )2/3 1/3

1/3

= (4π )1/3 (n B (1 + x)2/3 σ AB + n C σ BC )(3VC )2/3 1/3

(6.22)

1/3

= (4π )1/3 (n B x 2/3 σ BC + n C (1 + x)2/3 σ AC )(3VC )2/3

where x = VB /VC , Vi being the volume fraction of phase i; ni is the particle numbers of phase i in the system. Both the above models have been used extensively to predict the morphology of ternary or quaternary blends [7, 33, 35, 37, 94, 217, 221, 222, 228, 231, 233, 234]. For example, Hobbs et al. [37] found that the morphologies of ternary and quaternary blends they investigated are consistent with the calculated spreading coefficients for each polymer blend. The morphologies obtained for different HDPE/PP/PS and

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HDPE/PMMA/PS ternary blends agree with those predicted by the minimum free interfacial energy model [33]. Hemmati et al. [217] showed that the aforementioned two models predict almost the same morphology for PP/rubber (EPDM or EPR)/HDPE or PS (70/15/15 wt) ternary blends. The PP/rubber/HDPE blends show a core–shell morphology in which the HDPE core is encapsulated by the rubber shell. In PP/rubber/PS blends, the two minor phases remain dispersed separately in the PP matrix. Moreover, the results showed that interfacial interaction between phases is a major parameter that controls the phase structure in ternary blends. Melt viscosity ratio of components has no appreciable influence on the type of morphology and can affect the size of only dispersed phases, at least with the preparation mode used in their work. Virgilio et al. [234] also demonstrated that the prediction of the spreading coefficient model is consistent with the AFM micrographs of the HDPE/PS/PMMA blend. By changing the interfacial tension of a specific polymer pair in the ternary blend, it is possible to manipulate the change of the phase morphologies of multiphase blends from a classical matrix/dispersed phase system into a composite droplet system, or vice versa [33, 235]. Guo et al. [33] demonstrated that the phase morphology of a HDPE/PP/PS ternary system (with HDPE as the matrix) changes from an encapsulation type PP/PS to a separate droplet type by adding a small amount of compatibilizer to the blend. This is due to the fact that the reduction in the PS/HDPE interfacial tension reduces the spreading coefficient λPP/PS to a negative value. Legros et al. [235] showed the reverse effect for PBT/HDPE/EVA ternary blends. This blend demonstrates very little subinclusion behavior. However, after the addition of a selective compatibilizer for PBT/EVA, the blend forms a composite droplet morphology with EVA encapsulating HDPE. The in situ formation of PBT/EVA copolymer results in a reduction in the PBT/EVA interfacial tension, leading to a much more positive spreading coefficient λEVA/HDPE . For a PS/SBR/LDPE ternary blend, Luzinov et al. [34] observed a PS matrix, SBR shell, and SBR particles within the LDPE core. They suggest that this is due to the tendency of LDPE to envelop SBR at LDPE contents larger than the theoretical phase inversion point. In that case an impeded phase inversion process is not unexpected since the SBR copolymer possesses segments that are miscible with each component. The SBR acts as a classic compatibilizer interpenetrating each phase and acting as an adhesion promoter between the phases. Consistent observation was obtained by Reignier and Favis [94], who showed that it is possible to manipulate the internal structure of dispersed phase from small PMMA subinclusions dispersed in a larger PS particle to a PS/PMMA core–shell structure upon decreasing the PS/PMMA composition ratio in HDPE/PS/PMMA ternary blends (Figure 6.19). This core–shell structure is the result of an impeded phase inversion phenomena. From Figure 6.19 it can be observed that, whatever the PMMA content, it systematically forms a subinclusion structure in the PS phase, as expected by calculation of the spreading coefficients. For a high content of PS (Figure 6.19a), the phase morphology formed by the two minor components is constituted of distinct particles of PMMA dispersed in a larger and thus continuous phase of PS. These distinct particles coalesce and increase in size with increasing PMMA content (Figure 6.19b), as is observed in a binary blend. When sufficient PMMA is added to provoke binary blend phase inversion, a particular effect is observed. The PMMA subinclusions cannot go to the outside to become the continuous phase because of the opposing thermodynamic driving forces. Further increasing the PMMA content beyond the phase inversion leads to a complete core–shell structure with occasional ligaments spanning across the particle (Figure 6.19c) due to incomplete coalescence. This is a clear indication that thermodynamic forces are strong enough to prevent the phase inversion process. Further increasing the PMMA concentration leads to the appearance of PS particles of the shell material in the PMMA phase (Figure 6.19d). Here interfacial tensions between the phases are substantially higher and very little intermixing is expected between the phases, the thermodynamic driving forces are still strong enough to prevent phase inversion. Moreover, Reignier and Favis [94] showed that the composite droplet size reduces significantly for the 80:20 HDPE/(PS+PMMA) ternary blend with increasing the PS content. This is related to the presence of the PS shell at the HDPE/PMMA interface, which lowers the interfacial tension and also reduces the

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85% PS

Core/shell structure

65%

60%

35%

(a)

10% PS (d)

Impeded phase inversion

PMMA extracted

(b)

PMMA extracted

PS extracted

(c)

PMMA extracted

Figure 6.19 Evolution of dispersed phase internal structure as a function of PS content (vol% in dispersed phase) for 80:20 HDPE/(PS+PMMA) ternary blend. Reprinted from [94]. Copyright (2000) with permission from American Chemical Society.

matrix/dispersed phase viscosity ratio with increasing the outer shell thickness. These results also indicate that the composite droplet size resulting from breakup and coalescence is no longer influenced by the PS/PMMA ratio once a PS critical shell thickness of 0.2 μm is achieved. Above that thickness the composite droplet behaves in a fashion similar to a pure PS droplet. In a subsequent paper, Reignier and Favis [236] demonstrated that the PS/PMMA composite droplet exhibits pure PS droplet behavior at a critical volume fraction of encapsulating phase (PS:PMMA∼60:40). This critical volume fraction is independent of the overall dispersed phase concentration, shell thickness or dispersed phase size. Furthermore, the effect is observed even though the PMMA is significantly more viscous than the encapsulating PS phase. HDPE/PS/PMMA ternary blends have also been investigated by Guo et al. [33] and Reignier et al. [213, 233]. The results showed that for 80:20 (vol) HDPE/(PS + PMMA) blends, the system consists of a matrix of HDPE with PMMA/PS core/shell composite droplets, as predicted by the spreading coefficients. By carefully choosing the polymers and processing conditions, all polymers are perfectly segregated [233]. Control of the processing conditions also results in a nearly constant droplet size, while variation of the volume ratio of PS/PMMA allows for the obtention of PS shells ranging from 200 nm for 50:50 PS/PMMA composite droplets (based on the total dispersed phase volume, which is maintained constant at 20%) to theoretically 20–30 nm thick shells for 14:86 PS/PMMA composite droplets, which is in the range of a molecular layer of PS. Both aforementioned spreading coefficient and minimum free interfacial energy models are based on static interfacial tension. However, in melt-mixing of nonreactive polymer blends, predictions based on static interfacial tensions may be incorrect due to the difficulty in obtaining accurate interfacial tension data and the

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influence of shear flow on the delicate balance of interfacial tensions within a blend [212, 213]. Therefore, Reignier et al. [213] introduced a dynamic interfacial tension term (Eq. (6.13)), i.e. taking into account the elasticity of the polymer components, into the minimum free interfacial energy model (Eq. (6.22)), replacing the static interfacial tension. They obtained the following equations:     Ri Ri = 4π Ri2 σ B A + (N1,B − N1,A ) + 4π Ri2 σC A + (N1,C − N1,A ) B+C 6 6       Re Ri 2 2 (N1,B − N1,A ) + 4π Ri σC B + (N1,C − N1,B ) = 4π Re σ B A + Ai σi j B/C 6 6       Re Ri (N1,C − N1,A ) + 4π Ri2 σ BC + (N1,B − N1,C ) = 4π Re2 σC A + Ai σi j C/B 6 6



Ai σi j



(6.23)

where Re and Ri are the external and internal radius of the core–shell droplets, respectively; N 1 is the first normal stress difference for phases A, B, and C. The resultant conceptual model (Eq. (6.23)) can be called as a dynamic interfacial tension model. The results presented by Reignier et al. [213] strongly indicate that encapsulation phenomena in composite droplets is dominated by surface free energy considerations and that the dynamic interfacial tension needs to be taken into account. It was found that this dynamic interfacial tension model is able to reasonably explain the encapsulation behavior observed for three different blends composed of various PS and PMMA materials dispersed in an HDPE matrix. Valera et al. [228] studied the morphologies of PMMA/PP/PS blends with different compositions and compared the observed results with the predictions of aforementioned three phenomenological models. A mixture of separated PS droplets and core (PP)–shell (PMMA) morphologies is observed when PMMA is the matrix, a core (PMMA)–shell (PS) morphology is obtained when PP is the matrix, and separated dispersion is observed when PS is the matrix. The experimental observation corroborates the predictions of the free interfacial energy model only when PS is the matrix phase. The predictions of the spreading coefficient and dynamic interfacial tension models corroborate with most of the experimental results. However, none of the studied models is able to predict the present of pure PS droplets in the PMMA matrix. 6.3.4

Compatibilization

Controlling the phase behavior and morphology is crucial to determining the properties of polymer blends, which mainly rely on the interface between polymer components [237]. However, the inherent immiscibility of most polymer pairs and in a number of cases high interfacial tensions counteract a fine dispersion of the components during blending and result in coarsening phenomena such as coalescence (Section 6.2.3) and Ostwald ripening (also known as evaporation-condensation) [238]. Furthermore, most immiscible polymer blends suffer from a poor interfacial adhesion which causes inferior mechanical properties in the solid state [239]. Therefore, the interfacial modification, i.e. compatibilization, in immiscible polymer blends is a key issue for optimizing their morphologies and mechanical properties. Typically, compatibilization can be obtained by using block or graft copolymers, that is, compatibilizers, with segments which exhibit intermolecular attraction and/or chemical reactions with the blend components; these copolymers concentrate preferentially at the interface between the blend components. The compatibilizers can be pre-synthesized and then incorporated to a polymer blend in a separate step, i.e. physical compatibilization, or synthesized in situ during blending, by reactive compatibilization using suitably reactive monomers or polymers [15, 18, 193, 232, 240–242]. One of the disadvantages of the physical compatibilization is the tendency of the added copolymers to form micelles, which reduces the compatibilizing efficiency, increases the blend viscosity and may lessen the mechanical properties [22]. The reactive compatibilization

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is more complex than the physical compatibilization, but some studies (such as [103]) showed that the former achieves better compatibilization effects than the latter. Blends prepared by reactive compatibilization have thicker interphase than blends by physical compatibilization [22]. Reactive compatibilization has been shown to effectively stabilize the morphology of a number of polymer blends, such as ABS/PA6, PPO/PA66, PP/PA, PE/PA, PA/PS, PS/PP, PS/PMMA, PET/PC, PC/ABS, and so on [22, 103, 186, 193, 223, 240]. 6.3.4.1

Roles of Compatibilizers

There are three main roles of a compatibilizer in the blending process [22, 103, 169, 186, 223, 240, 243–245]. Firstly, it decreases the interfacial tension between the blend components and so retards the formation of the Rayleigh disturbances on the generated threads of dispersed phase. The lower the interfacial tension, the longer the deformation tension exceeds the interfacial tension, the longer the stretching of the thread proceeds, the smaller the diameter of the resulting thread becomes, and, consequently, the finer the size and dispersion of the generated dispersed phase. Secondly, it inhibits the coalescence process of the dispersed phase via steric stabilization during subsequent blending and is thus able to form and stabilize a finer morphology. Finally, it improves the interfacial adhesion in the solid state. This can be obtained either by ascertaining that an appropriate concentration of covalent bonds crosses the interface, by adding a compatibilizer that can act as an adhesive between two components, and/or by controlling the morphology, especially by inducing the phase co-continuity. The enhanced adhesion facilitates efficient stress transfer from one phase to the other phase and prevents cracks initiated at the interface from growth until the occurrence of catastrophic failure. Several investigations have shown that the incorporation of a compatibilizer to polymer blends results in the decrease of interfacial tension [22, 26, 187, 194, 205, 246, 247]. Mekhilef et al. [246] found that adding a triblock SEBS copolymer to the PS/HDPE blend prepared by TSE reduces the interfacial tension from 5.6 to 1.1 mN/m, which is significantly larger than the reduction in the corresponding dispersed phase particle size (×2). The interfacial tension in 75/25 LLDPE/PVC blend is reduced from 3.4 to 1.1 or 2.3 mN/m when using poly(isoprene-4-vinylpyridine) or hydroxyl-terminated PS as a compatibilizer [187]. P¨otschke et al. [247] used PE-g-MA as a compatibilizer in TPU/PE blends and found that the dispersed PE phase size decreases because of the reduction of the interfacial tension and viscosity ratio. Wu et al. [248] prepared PCL/PLA blends by melt-mixing using the block copolymers as compatibilizer and found that in the presence of copolymers, the interfacial tension reduces remarkably, accompanied by the thickening of the interface layer. As a result, an extra relaxation caused by Marangoni stresses is observed on the compatibilized blends. A dramatic reduction in the coalescence rate has been reported when a small amount of compatibilizer is added to polymer blends [103, 193, 194, 249]. For PS/EPR blend, Cigana et al. [249] displayed that measurable coalescence is initially observed at about 5 vol% EPR without a compatibilizer, whereas coalescence is suppressed up to 20 vol% EPR when the blend is compatibilized by a tapered poly(styrene-hydrogenated butadiene) diblock copolymer. For 70/30 PS/PMMA blends with symmetric P(S-b-MMA) diblock copolymers, Macosko et al. [193] found that preventing dynamic coalescence results in phase size reduction, while preventing static coalescence leads to morphology stability or compatibilization. They estimated that less than 5% of the interface needs to be covered to prevent dynamic coalescence, while about 20% is necessary to impart static stability. On the other hand, due to the fact that the droplet size decreases with compatibilizer, the number of droplets per volume unit increases and the separation distance between droplets decreases at a constant volume fraction of the dispersed phase. This results in an increase of flow-driven coalescence [189]. Some studies [21, 103, 122] suggested that compatibilizers promote morphological refinement much more by suppressing coalescence through stabilizing the interface during processing than by reducing the interfacial tension. Hu et al. [122] presented quantitative data for the coalescence efficiency in a linear flow by the visual observation of two equally sized drops. For very small concentrations of block copolymer, which causes negligible changes in interfacial tension, a drastic reduction of coalescence is observed.

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NON-COMPATIBILIZED

breakup/coalescence equilibrium

coalesce

e nc DROPLETS

comp

ati b

iliz

on ati

diffusion of block copolymer to new Interface

TIME

• D < 1 μm • coalescence inhibited COMPATIBILIZED

Figure 6.20 Schematic of morphology development during melt blending without and with a compatibilizer. Reprinted from [193]. Copyright (1996) with permission from American Chemical Society.

From the aforementioned results, it is possible to conclude that the effect of interfacial modification on the morphology predominates over that of composition and viscosity ratio of blends. Interfacial modification can reduce the size of the dispersed phase and to narrow its size distribution. Figure 6.20 intuitively illustrates the effect of a compatibilizer on the blend morphology development. If enough block copolymer can diffuse to the freshly generated interface, it should reduce the interfacial tension, permitting sheets to be drawn down thinner and prevent drop coalescence. Studies on the dependence of the dispersed phase size as a function of compatibilizer concentration have been undertaken for several polymer blends, such as PO/PA6 [250], PS/PA6, PS/HDPE, PP/PS, PS/elastomer, PET/PP, and LLDPE/PVC. The results demonstrated that the shape of the resulting curve is similar to that of the typical emulsification curve: an initial significant drop in dispersed phase size (especially the volume average diameter) at low concentrations of compatibilizer, followed by a levelling off to a plateau value at a certain critical concentration (Ccrit ), as shown in Figure 6.21. This Ccrit is believed to correspond to a point of saturation of the interface. Based both on emulsification curves mentioned above and interfacial tension measurements, the distinct role played by interfacial tension and by coalescence on the morphology of a polymer blend compatibilized by a given compatibilizer can be distinguished. Such a study was carried out by Lepers et al. [230]. They compatibilized the 90/10 and 99/1 PET/PP blends using a SEBS-g-MA, and generated the emulsification curves (Figure 6.21) and measured the interfacial tension. From Figure 6.21, it can be seen that without compatibilizer, the dispersed phase size is significantly higher for the 90/10 blend than for the 99/1 blend. This is due to the fact that significant coalescence occurs in the former, whereas coalescence is unlikely in the latter. In the 99/1 blend, the decrease of dispersed phase size is caused only by the decrease of the interfacial tension with the compatibilizer. It is clear for the 90/10 blend that both coalescence and interfacial tension are equally important in determining dispersed phase size reduction during compatibilization. It can be expected that if the dispersed phase concentration is to be increased to 20% or 30%, the decrease in particle size after

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4 Dispersed phase size (μm)

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2 Ccrit

Residual coalescence effect

1 Reduction due to interfacial tension 0 0.0

5.0 10.0 15.0 % Compatibilizer in the dispersed phase

20.0

Figure 6.21 Dispersed phase size as a function of compatibilizer concentration (based on the dispersed phase) for 99/1 and 90/10 PET/PP blends. The relative contribution of reduced coalescence and interfacial tension are shown. Reprinted from [230]. Copyright (1997) with permission from John Wiley & Sons.

compatibilization would be mainly due to a reduction in coalescence. Moreover, as indicated in Figure 6.21, the dispersed phase sizes at interfacial saturation for both blends are almost the same. This implies that for the 90/10 blend at saturation of the interface, almost all of the coalescence phenomena are suppressed by the addition of compatibilizer. Only a small residual coalescence effect remains for the 10% as compared to the 1% case at compatibilizer concentrations greater than Ccrit . Besides its concentration, the type of the compatibilizer also plays a significant role in determining the dispersed particle size [250]. Some studies were reported on the effect of compatibilization during the whole polymer blending process, including the initial stage and intermediate stage where major morphology development occurs (Section 6.2.1). The research undertaken by Cartier and Hu [251] showed that the morphology of in situ compatibilized PP/PA6 blends develops much rapidly than that of uncompatibilized counterparts. The final morphology is reached as soon as the phase transition is completed. This suggests that compatibilization also has an important effect on the initial breakup process which determines the final blend morphology to a great extent. This is probably due to the rapid creation of fresh interface at the early stage of mixing, which will be further analyzed in Section 6.3.4.3. Martin et al. [252] investigated the reactive processing of PBT/E-MA-GMA blends. It was shown that the morphology development can be subdivided in two successive steps (Figure 6.22). During initial mixing, the morphology evolution is essentially governed by the physical dispersion process, that is, breakup and coalescence. This is the ‘physics-controlled’ step (Zone I in Figure 6.22), in which the morphology evolution should be similar to that of the uncompatibilized blends. As the mixing time increases, the chemical reactions take place progressively, leading to a further refinement of the morphology and its subsequent stabilization (Zone II in Figure 6.22). In other words, the system shifts from ‘physics-controlled’ to ‘chemistrycontrolled’ steps. In this latter step, the morphology evolution is governed by the relative kinetics between

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Dn I when : Shear rate or/and PBT viscosity increases

II

when : Shear rate increases or/and T increases or/and PBT viscosity decreases Mixing time

when : T increases or/and Shear rate decreases

Figure 6.22 Schematic of phase size reduction in reactive processing. I: ‘physics-controlled’ step; II: ‘chemistrycontrolled’ step. Reprinted from [252]. Copyright (2004) with permission from John Wiley & Sons.

the compatibilization reaction and the crosslinking of the rubber phase. The shift between the two steps is progressive and intimately dependent on the processing conditions (as shown in Figure 6.22) and reactive kinetics. Recently, Li and Sundararaj [18] investigated the mechanism of how compatibilization affects the morphology development during the initial stage of polymer blending in the perspective of hydrodynamic forces. Model experiments of 80/20 PMMA/amorphous PA (aPA), PS/aPA, and PS-g-MA/aPA blends were performed. Force analysis describing the phase deformation during blending indicated that compatibilization enhances the dispersion of blends by reducing the slip at the interface between the polymer phases. Polymer segments of block copolymer formed by in situ reaction penetrate into both the matrix and dispersed phases and entangle with the bulk polymer chains, thus improving the adhesion between the matrix and dispersed phases. Therefore, the shear stress from the matrix phase to the dispersed phase can be transferred more effectively. 6.3.4.2

Mechanisms of Deformation and Breakup of Compatibilized Droplet

In the presence of a compatibilizer, the mechanisms of deformation and breakup of dispersed phase change substantially. The drop deformation is not only controlled by the ηr and the Ca, but also by two additional parameters: the surface P´eclet number (Pes ) and interfacial tension ratio (I r ) [253]. Pes is defined as the ratio between the surface convective flux, JC = σ γ˙ a, that promotes the concentration gradient, and the surface diffusion flux, J D = D S σ/a, that tends to restore homogeneous concentration distribution of the compatibilizer molecules along the drop surface: Pes = a 2 γ˙ /D S

(6.24)

where a is the particle diameter, σ is the interfacial tension under quiescent conditions, and DS is the surface diffusion coefficient of the compatibilizer molecules. I r is given by: Ir = σ (x, t)/σ0

(6.25)

where σ 0 is the interfacial tension in the absence of the compatibilizer (clean surface) and σ (x,t) is the interfacial tension in the presence of the compatibilizer that changes in time along the drop surface. In turn,

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such spatial variation in the interfacial tension changes the capillary number that becomes: ∗ Cacrit = Cacrit /Ir

(6.26)

∗ where Cacrit and Cacrit are the critical capillary numbers of compatibilized and uncompatibilized drops, respectively. As mentioned in Section 6.3.4.1, the presence of a compatibilizer decreases the interfacial ∗ and consequently changes the tension. One can see from Eqs. (6.25) and (6.26) that this increases the Cacrit drop deformation [122, 241, 245]. If the deformation is large, the compatibilizer concentration distribution along the surface becomes nonuniform, which results in a nonuniform drop deformation. Sailer and Handge [239] investigated the influence of reactive compatibilization on the deformation of the dispersed phase in elongated PA6/SAN blends. The stretch ratio of the reactively compatibilized SAN drops in the PA6 matrix is larger compared to that of the drops in the neat blends and increases with SAN-gMA concentration. Figure 6.23 shows the AFM micrographs of the elongated samples taken parallel to the stretching direction. The initially spherical SAN drops are stretched into an ellipsoidal shape. Increasing the SAN-g-MA concentration from 0 to 3.4 wt% leads to a larger stretch ratio for the drops. This can be explained by a decreased interfacial tension and by the flow-induced accumulation of compatibilizer molecules at the drop tips as shown in Ref [122]. However, a further increase to 4.8 wt% SAN-g-MA does not seem to change the stretch ratio of the drops any more. This indicates that already at 3.4 wt% SAN-g-MA the interface is saturated with multigrafted SAN-g-MA chains. Some researcheres [15, 253, 254] studied the effect of compatibilizer on polymer drop deformation and breakup during shearing using a visualization approach. Levitt and Macosko [254] sheared PP or PMMA drops in PS matrix and observed that the addition of compatibilizer results in long, sheet-like droplets, i.e. the transition from affine to fiber-like deformation is delayed. In some cases, compatibilized drops even display widening, i.e. the dimension along the vorticity direction is larger than the original droplet diameter. This may be attributed to a combination of reduction in interfacial tension and interfacial slip as well as a gradient in interfacial tension (Figure 6.8b). Lin et al. [15] found that compared to uncompatibilized PE/PS blends, the presence of premade diblock copolymer does not affect initial breakup of PS drop sheared in a PE matrix significantly. When premade copolymer is added inside the PS drop, the drop interfacial concentration is lower than the saturation coverage (as shown in Figure 6.24a). For a reactive system, when a PS oxazoline (PSOX) drop is sheared in a PE matrix with PE-g-MA, surface erosion and vorticity elongation and breakup of drop occur at a higher shear rate than the nonreactive system. This may be due to the fact that the crosslinking reaction between PSOX and PE-g-MA takes place gradually at the interface, forming a small amount of crosslinked copolymer (as shown in Figure 6.24b) that prevents breakup. In the presence of premade or reactively formed copolymer, a tiny cylindrical tip develops in the vorticity direction and then ruptures from the elongated mother drop (Figure 6.24). The formation of this tiny tip with high interfacial curvature stabilizes the mother drop. Very recently, Abbassi-Sourki et al. [253] visualized the effect of a PDMS-PIB block copolymer on the deformation and breakup behavior of Newtonian PDMS drops in a Newtonian PIB matrix. Uncompatibilized PDMS drops deform symmetrically and form a waist prior to breakup into an odd number of smaller droplets. By contrast, the compatibilized drops deform asymmetrically, sharpen at the ends and finally go through tip-dropping or end-splitting. Deformation of the lower viscosity ratio drops is accompanied by tip-stretching and tip-streaming (as shown in Figure 6.6). The authors calculated the Cacrit from experimental breakup data and the results are shown in Figure 6.25. The Cacrit of uncompatibilized drops follow the same trend as the experimental results of Grace [134] fitted by Eq. (6.9), with a minimum in Cacrit for ηr lying between 0.41 and 1.1, but are smaller than the ones reported by Grace for all viscosity ratios. The Cacrit for compatibilized drops is approximately one order of magnitude higher than that of uncompatibilized drops. This is due to the fact that in the presence of the block copolymer the interfacial tension decreases

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Figure 6.23 Atomic force micrographs of 70/30 PA6/SAN blends with (a) 0, (b) 0.6, (c) 2.0, (d) 3.4, and (e) 4.8 wt% SAN-g-MA. The samples were elongated to the maximum Hencky strain of 1.8 with Hencky strain rate of 0.3 s−1 at 240 ◦ C. The direction of stretching is indicated by an arrow. Reprinted from [239]. Copyright (2008) with permission from American Chemical Society.

more than the critical shear rate. However, the drops with the highest and lowest viscosity present a lesser increase in the ratio of (Cacompat /Cauncompat )crit . Some contrary results, that is, the compatibilized drops are more stable against breakup, were reported. Velankar and co-workers [241, 255] studied model blends (PIB/PDMS, with a ηr of about 1) compatibilized with 0–10% PIB-PDMS diblock copolymer. They found that for blends with 1% PIB, the PIB droplet shows pointed ends in the velocity direction during shearing upon addition of 2% copolymer. The authors ascribed this to the accumulation of copolymer at the droplet tips and suggested that this accumulation results in a lower local interfacial stress, giving a higher local curvature to balance the pressure jump across the interface. When adding 10% copolymer, the PIB droplet almost does not deform. Moreover, they measured the Ca during coalescence for blend with 10% PIB. The addition of the copolymer increases the steady shear Ca of droplets to values well above the Cacrit for breakup of uncompatibilized droplets. This suggests that the hydrodynamic

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Copolymer premixed with the drop and the drop interfacial Tiny tip develops concentration below saturation

Tiny tip breaks off

(b)

Small amount of crosslinked product formed at the interface

Tiny tip develops

Breakup of the tiny tip, drop erosion and vorticity breakup at higher shear rate

Figure 6.24 Schematic of effects of (a) premade copolymer and (b) reactively formed copolymer on polymer droplet breakup. The flow and vorticity directions are horizontal and vertical, respectively. Reprinted from [15]. Copyright (2005) with permission from American Chemical Society.

stress required to break compatibilized droplets is considerably higher than that expected on the basis of their interfacial tension. That is, compatibilized droplets show a significant decrease in the deformability in shear flow compared with uncompatibilized ones. This, as well as the aforementioned undeformed droplet upon 10% copolymer, may be explained as follows. The deformation of the compatibilized droplet creates gradients in the compatibilizer concentration on its surface. This results in substantial gradients in interfacial tension (Marangoni stresses) along the interface, which immobilize the interface at high surface coverage and so stabilize the droplets. 1000 Eq. (3) Eq. (21)

100 Compat. drops Cacrit

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0.1 1e–6

Uncompat. drops

1e–5

1e–4

1e–3

1e–2

1e–1

1e+0

1e+1

viscosity ratio

Figure 6.25 Evolution of Cacrit for uncompatibilized and compatibilized droplets. Reprinted from [253]. Copyright (2009) with permission from Elsevier.

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The above presumption was later verified by Jeon and Macosko [256], who visualized the fluorescent block copolymer distribution on the surface of a PMMA droplet in a PS matrix. Confocal microscopy revealed that the block copolymer distributes uniformly on the droplet surface before deformation. However, shear deformation results in higher concentration at the droplet edges and tips, which is believed to be caused by the convection of copolymer induced by shear flow. 6.3.4.3

Effect of Processing on Compatibilizing Efficiency

Flow in melt mixers tremendously accelerates the compatibilizing efficiency. This may be caused by the rapid generation of interfacial area during early mixing and by the convection. As illustrated in Figures 6.3 and 6.20, initial morphology was found to be ribbons or sheets developing to a lacelike structure [30, 85, 86, 89, 101], thus the extent of deformation is greatest at this initial stage, where the interfacial area increases by about 1000 times. In addition, flow may change the concentration profile of functional groups at the interface and/or increase the collision probability [244]. It has been shown that the mixing procedure (one-step and two-step mixing) may significantly affect the capacity of the compatibilizer to migrate to the interface, especially for blends in which the compatibilizer has a higher affinity with one of the components. Generally, when the compatibilizer is initially blended with the component for which it has more affinity and the resulting blend is subsequently compounded with the one for which it has less affinity, the migration of compatibilizer to the interface is facilitated and so finer dispersions are obtained [169]. This, in turn, influences the emulsification and subsequent morphology of the blends. Cimmino et al. [257] prepared ternary blends by adding to PA6 an EPM and the same EPM functionalized with maleic anhydride (EPM-g-MA). One-step and two-step mixing procedures were used. In the former, three components are simultaneously introduced into the mixer; whereas in the latter, two rubbers (EPM and EPM-g-MA) are separately premixed before the final mixing with PA6. It was shown that an EPM-g-PA6 graft copolymer is formed, which acts as a compatibilizer between the dispersed rubbery phase and the PA6 matrix. The blends prepared by the two-step mixing exhibit a very fine morphology, whereas the ones prepared by one-step mixing show a coarse morphology. Recent years have seen increasing efforts towards one-step reactive blending [242, 258] in which both the functionalization and the reactive blending steps are executed in the same extrusion process. It was demonstrated by Barangi et al. [242] that the grafting or compatibilizing efficiency achieved by the one-step procedure is greater than that obtained by a two-step procedure at a similar grafting level. The one-step compatibilized blends thus exhibit stronger interfacial interaction. Coltelli et al. [258] also showed that the one-step mixing seems promising to obtain toughened PA-based blends. 6.3.4.4

Compatibilization of Ternary and Quaternary Blends

Horiuchi et al. [231] studied the morphology development of immiscible ternary blends through an interfacial compatibilization reaction between the matrix and one dispersed phase by using SEBS-g-MA or PS-g-MA as the compatibilizer for ternary PA6/PC/SEBS and PA6/PC/PS blends, respectively. The added compatibilizers react with the amine end groups of PA6 matrix at the interface during melt-mixing. This interfacial reaction results in a reduction of the interfacial tension and a fine dispersion of the compatibilizers in the PA6 matrix at about 100 nm in diameter, and at the same time induces the change of the formation of the two dispersed phases from ‘stack formation’, where the two dispersed polymers are stuck together (as shown in Figure 6.18), to ‘capsule formation’, where the PC domains are encapsulated by the other phase. Moreover, as shown in Figure 6.26, in the PA6/PC/SEBS blend, the encapsulation by SEBS onto the PC domains gradually becomes complete as the ratio of SEBS-g-MA to unmodified SEBS increases, and complete encapsulation can be achieved only by the use of SEBS-g-MA alone; whereas in the PA6/PC/PS blend, the complete

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SEBS+SEBS-gMA

SEBS-gMA

SEBS PC PC

PA6

PC

PA6

PA6

Interfacial reactivity PS+PS-gMA

PS-gMA

PS PC PC

PA6

PC

PA6

PA6

Figure 6.26 Schematic of tendency of morphology development for (a) PA6/PC/SEBS and (b) PA6/PC/PS blends by varying the ratio of unmodified polymer to its maleinated polymer. Reprinted from [231]. Copyright (1997) with permission from American Chemical Society.

encapsulation by PS onto PC can be obtained even when both PS and PS-g-MA are used. This suggests that low interfacial reactivity is enough to give complete encapsulation in the PA6/PC/PS blend, whereas higher interfacial reactivity is required to achieve the complete encapsulation in the PA6/PC/SEBS blend. Tomova and Radusch [259] evaluated the phase interactions in ternary PA6/PA66/elastomer blends with two different functionalized elastomers, that is, EO-g-MA and EPDM-g-MA. The atomic force microscopy study showed that the blends exhibit phase-in-phase and ‘quasi’ phase-in-phase morphologies when the former and latter copolymers are used elastomer phases, respectively. Incorporation of PA inside the elastomer particles was observed in the former case presumably due to the disfunctionality of PA66. de Freitas et al. [260] studied the morphology evolution of PP/PS/PMMA ternary blends, with or without compatibilizer (PP-g-PS, SEBS, or PP-g-PMMA). All the studied blends present a core–shell morphology with PS as shell and PMMA as core. The addition of PP-g-PS or SEBS decreases the core–shell diameter of the composite droplets and the addition of PP-g-PMMA does not seem to have any effect on the diameter of PMMA. Moreover, when PP-g-PMMA is added to the PP/PS blend, some PMMA migrates from the core of the PS droplet to the PS/PP interface. The difference observed in the final morphologies relies on the number of PMMA core particles present in the PS shell. The morphologies predicted using the spreading coefficient and minimum free energy models (Eqs. (6.20) and (6.21)) corroborate the equilibrium morphologies experimentally observed for the ternary blends, except when PP-g-PMMA is added to the blend. Very recently, Virgilio et al. [261] tried to prepare ultraporous PLLA scaffolds by melt-processing quaternary EPDM/PCL/PS/PLLA blends and found that the PLLA scaffolds modified with the PS-b-PLLA diblock copolymer present themselves as fully interconnected porous networks with asymmetric channel walls.

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Moreover, reactive compatibilization can provide for a degree of control over morphology development in ternary or quaternary polymer blends, via manipulation of the interfacial energies within the blends, which allows the formation of a composite droplet morphology during melt blending via encapsulation of one dispersed phase by another [37, 33, 213, 231]. 6.3.4.5

Compatibilization of Nanocomposites

In contrast to polar polymers, like PAs, that can effectively exfoliate organically modified clays using conventional melt processing techniques, for nonpolar polymers, such as the most widely-used polyolefins, PE or PP, well exfoliated and homogeneous dispersion of the silicate layer at the nanometer level appear to be more difficult [50, 262, 263]. Consequently, the matrix modification with polar moieties is necessary prior to modified clay introduction to achieve nanometric dispersion of the clay. Initial attempts to create nonpolar polymer/clay nanocomposites by simple melt-mixing were based on the introduction of a modified oligomer to mediate the polarity between the clay surface and the polymer [264]. One typical example is PP/clay nanocomposite: the nanodispersion of the organoclay is achieved using PP-o-g-MA as a compatibilizer. It has been found that polyolefin oligomers with polar functionality can be intercalated into the organophilic clay galleries during melt-blending and these phenomena are the key to achieve the dispersion of exfoliated silicates in polyolefins [79, 265, 266]. Using this approach, Usuki et al. [265] first prepared PP nanocomposites. The functional oligomer is intercalated between the layers of clays, and then the oligomer/clays is melt-mixed with PP to obtain the nanocomposite with intercalated structure. Kawasumi et al. [79] successfully prepared PP-clay hybrids by simple melt-mixing using a PP-o-g-MA as a compatibilizer. The miscibility of the PP-o-g-MA with PP affects the dispersibility of the clay in the hybrids. Almost completely dispersed hybrids are obtained in the case where the used PP-o-g-MA is miscible with PP (as shown in Figure 6.9). Moreover, only the PP-o-g-MA can intercalate into the silicate layers at the first stage of the mixing process. Hasegawa et al. [266] blended the PP with organophilic clay that is intercalated with a PP-o-g-MA, which is also used as a compatibilizer. In these prepared PP-clay hybrids, some particles of silicate layers are dispersed at the nanometer level. The particles become smaller and are dispersed more uniformly, as the ratio of the PP-o-g-MA to the clay is increased. It becomes apparent that the PP-o-g-MA improves the dispersibility of the clays in the hybrids. The most promising strategy at the present time is to add a small amount of a MA-grafted polyolefin that is miscible with the base polyolefin [50]. It is believed that the polar character of the MA has an affinity for the clay materials, such that the MA-grafted polyolefin can serve as a ‘compatibilizer’ between the polymer matrix and the filler. If the compatibility is sufficient to disperse the clay silicate layers at the molecular level, exfoliated-type polyolefin/clay nanocomposites can be easily prepared by melt-mixing. This strategy has been widely used for PP- and PE-based nanocomposites. Hasegawa et al. [80] prepared polyolefin-clay hybrids by melt-blending of MA-modified polyolefins (including PP-g-MA, PE-g-MA, EPR-g-MA) and organophilic clay. They described the clay dispersion process as follows. The modified polyolefins are first intercalated into the organophilic clay galleries during melt blending to increase its interlayer spacing while the interaction of the layers is weakened. The driving force of the intercalation originates from the strong hydrogen bonding between the MA groups and the polar clay surfaces. Then the layered structures of the intercalation compound are disordered by shear force during blending and the clay silicate layers are finally exfoliated and homogeneously dispersed at the monolayers. In an extended study by Hasegawa and Usuki [267] on the silicate layer exfoliation process in MA-grafted polyolefins, it was found that molten PP-g-MA continuously intercalates into the clay galleries and the silicate layers exfoliate spontaneously without shear. However, the silicate layers do not disperse into the PP-g-MA matrix until a TSE is used, that is, shear is induced. Under controlled conditions of shear history and temperature, Galgali et al. [268] studied the flow-induced clay orientation in extruded tape samples of

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syndiotactic PP nanoclay composites. They found that the clay orientation increases with increasing shear rate in compatibilized nanocomposites, but is not influenced by the shear rate in uncompatibilized hybrids. Many kinds of polyolefin/clay nanocomposites using MA-modified polyolefins have been prepared. Nam et al. [269] mixed PP-g-MA (0.2 wt% MA) and different amounts of organophilic clay in a TSE to prepare PP/clay nanocomposites. They obtained nearly exfoliated structures upon adding 2 wt% clay. However, the addition of 4 and 7.5 wt% clay results in disordered intercalated and ordered intercalated structures, respectively. L´opez-Quintanilla et al. [270] analyzed the effects of three compatibilizers (GMA, AA, and MA) and incorporation method on the clay dispersion in PP-clay nanocomposites. Results showed that clay dispersion and interfacial adhesion are greatly affected by the kind of matrix modification. PP-g-GMA and PP-g-MA are better compatibilizers than PP-g-AA. Moreover, better dispersion and exfoliation for the nanoclays are obtained when using two-step mixing than one-step mixing conditions. Via melt compounding, Kato et al. [271] prepared PE-clay hybrids by using PE-g-MA, in which the silicate layers are completely exfoliated and dispersed on a nanometer level in the PE matrix. However, nonmaleated PE does not exfoliate clay under any conditions. By adding MA directly as a compatibilizer during melt blending, Zhang and Wilkie [272] obtained PE-organoclay nanocomposites with a mixed immiscible-intercalated structure. The authors presumed that there is some reaction that occurs between the MA and the PE during the high temperature blending, leading to the formation of a graft copolymer in which MA units are attached to the PE chain. In addition to MA and PP-g-MA, EVA copolymer has also been used as a compatibilizer to prepare PE-based nanocomposites. For example, Zanetti and Costa [262] prepared composites with different PE/EVA ratios and 5 wt% organoclay by melt-mixing. No interaction is obtained by mixing the molten PE with the organoclay without a compatibilizer. Adding 1 wt% of EVA is enough to promote intercalation whereas increasing its amount to 5 wt% the interaction between the polymer matrix and the organoclay increases. Chiu et al. [273] used two maleated polyolefins (PE-g-MA and PP-g-MA) separately as a compatibilizer to prepare PP/HDPE blend nanocomposites. The results showed that the organoclay is intercalated and/or partially exfoliated within the blend upon adding PE-g-MA or PP-g-MA and the latter facilitates the dispersibility of the organoclay to a better degree. Compatibilization is also a critical issue for preparing nanocomposites based on other polymers, such as PS. As was demonstrated by Hasegawa et al. [197] the layered structures of silicates are observed at about 500–1000 nm in the PS/organophilic clay composites prepared using a TSE. So they introduced 46.5–93 wt% PSoz, being a compatibilizer between PS and clay, into the PS. In so-prepared nanocomposites by melt blending, silicate layers of the clay are delaminated and dispersed to the monolayer level (Figure 6.10), due to the existence of oxazoline groups. Using in situ reactive blending, Wang and Wilkie [274] prepared PS/clay nanocomposites with both the organically modified and the pristine inorganic clays and found that MA increases the possibility of nanocomposite formation. 6.3.5

Composition

During the changing of the composition, binary polymer blends generally exhibit two types of morphologies, i.e. the matrix-dispersed structure and the co-continuous structure. Usually, in the low composition range, one of the polymers is dispersed in a continuous phase of the other [275, 276]. Above a certain composition, phase inversion may occur and either co-continuous or fibrillar structures are obtained. Many researchers have reported that the size of the dispersed phase increases and its size distribution broadens significantly with the increase of the dispersed phase contents [103, 181, 183, 184, 247, 275, 277]. This is usually caused by coalescence. As noted in Section 6.2, the morphology development is governed by two competitive processes, i.e. breakup and coalescence of the dispersed phases. Compared with the breakup, the coalescence is strongly influenced by the dispersed phase contents [103, 184], which was observed at contents lower than 1% [184]. Actually, at higher dispersed phase contents, the final morphology results from

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a competition between breakup and coalescence, whereas, at low enough contents, breakup is the dominant effect that dictates the lower limit of dispersed phase size [240]. Some studies, such as [277], showed that at high dispersed phase fractions, dispersed phase sizes of blends with high values of ηr are smaller than those of blends with low values of ηr . This can be explained based on the fact that at high values of ηr , the overall viscosity increases as increasing the fraction of the dispersed phase. Some experimental results, contrary to the generally accepted view regarding the dispersed phase fraction effect, were also reported. For polymer blends with a values of ηr higher than unity, the particle size decreases with concentration of the dispersed phase up to a certain value, such as 20 wt% [181, 278]. This occurs because the total deformation of the dispersed phase, before breakup, increases as its content increases, and coalescence is hindered. Therefore, it can be expected that in industrial mixers, the size of the dispersed phase should decrease as its volume fraction increases, if coalescence is suppressed. 6.3.5.1

Prediction Models

Several models have been proposed for predicting the diameter of the dispersed phase affected by the blend composition. Serpe et al. [279] modified Wu’s equation (Eq. (6.10)) by substituting the matrix viscosity by the blend viscosity (ηb ) and by introducing a term that considers the blend composition and, thus, coalescence effects. The empirical relationship proposed by them is as follows: 4σ Dn = ηb γ˙ F(φ)



ηd ηb

±0.84

F(φ) = 1 − (4φd φm )0.84

(6.27) (6.28)

where φ d and φ m are the volume fractions of the dispersed phase and the matrix, respectively. Considering the competing processes of breakup and coalescence of the dispersed phase, several other researchers [183, 191] proposed the equations (see Section 6.2.3) for predicting its equilibrium diameter. Using the volume concentration ratio c φ (ratio between the volume component of the dispersed phase and the volume component of the continuous phase) and the ηr , it is possible to estimate the expected phase morphology. If c φ /ηr < 1, a matrix-dispersed morphology is possibly assumed; a co-continuous morphology is generally obtained for c φ /ηr = 1; for c φ /ηr >1.2, a phase inversion is assumed [90]. 6.3.5.2

Ternary Blends

Over the last years, increasing attention has been paid to the effect of composition on the morphology of ternary polymer blends. Luzinov et al. [34] investigated the morphology of ternary PS/SBR/LDPE blends at a constant content (75 wt%) of the PS matrix while changing the weight ratio of the two minor components (LDPE to SBR). They found that the size and ultimate structure of the dispersed phases of the ternary blends can be controlled by this weight ratio. Upon increasing this weight ratio, the dispersed phase changes from a multicore structure to a LDPE/SBR core–shell morphology. The size of the LDPE subphase in the mixed dispersed phase increases sharply at a LDPE content that corresponds to phase inversion in the parent SBR/LDPE binary blends. In a subsequent study of Luzinov et al. [7], the HDPE/PS weight ratio in ternary PS/SBR/HDPE blends is changed at constant SBR content (25 wt%). Depending on this ratio, two main types of phase morphologies, i.e. dispersed phases with a core–shell structure with SBR forming the shell and three co-continuous phases, are observed. At low HDPE (or PS) content, HDPE (or PS) is encapsulated by an SBR shell and dispersed in the PS (or HDPE) matrix (Figure 6.27a and c). When increasing the HDPE (or PS) content, the particles formed by this component are much larger and much irregularly shaped but still

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Figure 6.27 TEM micrographs for PS/SBR/HDPE blends with HDPE/PS composition (wt) of (a) 20/80, (b) 50/50, and (c) 80/20. Reprinted from [7]. Copyright (2000) with permission from Elsevier.

coated by an SBR layer. At HDPE/PS ratios of 40/60, 50/50 (Figure 6.27b) and 60/40, three co-continuous phases are formed, SBR tending to localize at the boundary between the co-continuous HDPE and PS phases. In addition to co-continuous phases, HDPE and PS particles enveloped by SBR also coexist. Core–shell morphology for the minor phase is observed at quite different HDPE/PS ratios, independent of the polymer that forms the matrix, either more viscous and elastic HDPE or less viscous and elastic PS. Observation of co-continuity region (triple-phase continuity) complies with the minimization of the total interfacial free energy for the system. Figure 6.28 illustrates the dependence of the average diameters of the HDPE (or PS)

6

6 PE matrix

PS matrix

5

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Co-continious morphology

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20

40

60

80

100 PE

PE/PS ratio (wt%)

Figure 6.28 Effect of HDPE/PS weight ratio on the number () and volume () average diameters of cores in PS/SBR/HDPE ternary blends. Reprinted from [7]. Copyright (2000) with permission from Elsevier.

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cores on the HDPE/PS weight ratio in the core–shell morphology. The core diameters expectedly increase with the content of the core-forming polymer (PS or HDPE). Hemmati et al. [280] investigated the effect of composition of components on morphology of PP/EPDM/HDPE and PP/EPDM/PS ternary blends, in which the PP is the matrix. The composition affects only the size of each dispersed phase, whether it forms core or shell or disperses separately, and has no appreciable effect on the type of morphology. The former blends show a core–shell-type of morphology in which EPDM encapsulates HDPE. The size of the HDPE core and dispersed phase as a whole is directly related to HDPE content. Increasing EPDM content at constant amount of HDPE increases the dispersed phase size, while reducing HDPE core size; whereas in the latter blends, EPDM and PS form separate phases. The EPDM particles are generally smaller than PS particles because of their lower interfacial tension with PP. Wilkinson et al. [232] prepared ternary PP/PA6/(mixed SEBS/SEBS-g-MA) (70/15/15, wt) blends with different SEBS/SEBS-g-MA ratios via melt blending. It was demonstrated that changing this ratio generates very different dispersed-phase morphologies. The progressive replacement of SEBS with reactive SEBS-gMA increases the degree of interfacial reaction between the succinic anhydride groups of the SEBS-g-MA and the terminal amino groups of the PA6, thus significantly reducing the PA6-SEBS interfacial tension and providing a driving force for encapsulation of the PA6 by the SEBS. Consequently, the dispersedphase morphology was observed to transform from two separate phases (100/0) to acorn-type composite droplets (Figure 6.18) (75/25), then to predominantly individual core–shell particles (50/50) and finally to large agglomerates of these core–shell particles (25/75 and 0/100). The transformation in dispersed-phase morphology produces a range of PP ternary blends with much superior ductility and impact-strength compared to the base PP, but at the cost of significant reductions in modulus and yield stress. For HDPE/PS/PMMA blends, Tchomakov et al. [38] found that the PS encapsulates the PMMA to form composite droplets within the PE matrix. Relatively stable morphology was obtained with the dispersed phase volume fraction of 1:1 over a wide range of processing conditions, polymer feeding sequence, minor phase viscosity and minor phase concentration. Reignier et al. [233] examined the issue of core–shell morphology formation within the dispersed phase for composite droplet polymer-blend systems composed of an HDPE matrix, PS shell, and different molecular weights of PMMA core material. Figure 6.29 shows the evolution of the PS shell formation for the 80/20 HDPE/(PS/PMMA) ternary blends with increasing PS content. It was shown that changing the molecular weight of the core has a dramatic effect on the PS-PMMA structure within the composite droplet. Moreover, the PS/LMW-PMMA system forms a complete PS shell structure at a significantly lower PS composition than the HMW-PMMA case. It is possible to control the shell thickness of the core–shell particles through the minor phase composition ratio without substantially affecting the composite droplet particle size. It was observed by Ha et al. [216] that in PP/mPE/HDPE ternary blends, the extent of encapsulation of HDPE is increased with increasing mPE concentration and the requirement for complete encapsulation seems to be >83 wt% mPE in mPE/HDPE binary blends. For PC/ASA/SAN blends with different compositions, Han et al. [281] showed that SAN and ASA contents have great influence on the morphology and properties. The ternary blends with a small amount SAN (such as 5 wt%) show good dispersed morphology and better integrated properties. Besides the original material properties and the phase morphology, the interphase (thickness, composition, and structure) between the different phases is also a key parameter determining the properties of immiscible polymer blends [234]. Some work has been done in immiscible melt-processed polymer blends to quantitatively characterize the interphase, or transition region, between the different phases in terms of thickness, structure, and composition. Using focused ion beam preparation followed by tapping mode atomic force microscopy, Virgilio et al. [234] demonstrated an approach for the analysis of model interphases in 80/20 HDPE/(PS/PMMA) ternary blends with core–shell morphology. The authors found that as decreasing the PS content, the PS shell thickness decreases, while the average diameter of composite droplets does not change significantly with PMMA/PS ratio. This effect was also observed by Reignier and Favis [233]. Recently, Zhang

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Figure 6.29 Evolution of shell formation process with increasing PS content (vol% based on the dispersed phase) for 80/20 HDPE/(PS/PMMA) ternary blends with (a) low- and (b) high-molecular weight cores. Black and gray phases represent the PS and the PMMA components, respectively. Reprinted from [233]. Copyright (2003) with permission from American Chemical Society.

et al. [282] investigated the development of ternary percolated co-continuous systems in HDPE/PS/PMMA blends at different compositions based on the co-continuity region found for HDPE/PMMA binary blends. HDPE and PMMA form two continuous networks, while the PS forms a continuous sheath structure at the HDPE/PMMA interface in the system. Results showed that a continuous shell structure of PS is situated at the co-continuous HDPE/PMMA interface through combineing a co-continuous structure with a composite droplet type of morphology. The presence of a PS layer at the HDPE/PMMA interface results in a dramatic decrease in the percolation threshold volume fraction of the encapsulating PS component. As little as 3% PS results in a PS phase continuity of about 70% (as shown in Figure 6.30), a very high level of continuity for such a small volume fraction of PS. 6.3.6

Nanoparticles

Over the last decade, more research attention has been focused on the influence of nanoparticles, especially organoclays, on the morphology of immiscible polymer blends, such as PE/PA6 [62, 82, 283], PP/PA6 [72, 84], LDPE/PA11 [284], PS/PP [67], PS/PMMA [69, 237], PE/PBT [74, 75, 285], PP/elastomer [83, 286, 287], PP/PET [288], PA6/ABS [71], PA6/EPR [61, 64], PA6/EPDM-g-MA [65], PPO/PA6 [70], PEMA/PS [73], PMMA/EVA [237], PC/SAN [237], PC/PMMA [66], PP/PBSA [68], and so on. Most results showed that, besides the reinforcement effect on polymer matrices, nanoparticles also decrease the size of dispersed

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(b)

(a)

HDPE PS PMMA PMMA

PS PS

HDPE

HDPE

Figure 6.30 Schematic of HDPE/PS/PMMA multiple percolated structure: (a) three-dimensional view of blend structure, showing thin PS layer at HDPE/PMMA interface; (b) section showing relative position of different phases. Reprinted from [282]. Copyright (2007) with permission from American Chemical Society.

phase obviously. Three possible mechanisms have been proposed regarding the phase size reduction of blends upon the addition of clays: (1) the compatibilizing effect between blend components; (2) the increase in the matrix phase viscosity; and (3) the effective suppression of the coalescence for dispersed domains. 6.3.6.1

Compatibilizing Effect

Some researchers confirmed that organoclays play a role as a compatibilizer for immiscible polymer blends [66, 67, 69, 73, 237, 283, 287]. There are two possible mechanisms for the organoclay compatibilization. First, the organic modifier (intercalant) of clay is miscible with both blend components, thus the overall freeenergy of mixing becomes negative and thermodynamically driven compatibility appears to occur between the immiscible components. Then, molecular chains of both components can diffuse into the clay galleries, resulting in a uniform dispersion of clay in the matrix phase or mainly at the interface between two phases. Secondly, the stabilizing energy gain originates from the adsorption of polymer components on the solid surface. In the case of organoclay or carbon nanotube, the energy gain becomes larger due to their large surface area per unit weight [285]. Voulgaris and Petridis [73] added 5–27% organoclay to the 75/25 PEMA/PS blend and found that the organoclay adsorbs selectively the PEMA chains and creates intercalated structures. The size of the PS domains decreases with increase in the clay concentration, giving a typical emulsification curve. So the organoclay exhibits an emulsifying action. Mehrabzadeh and Kamal [283] showed that adding 5 wt% organoclay to 80/20 HDPE/PA6 blend reduces the size of PA6 particle, and changes its morphology from spherical to laminar shape. Gelfer et al. [69] demonstrated that although immiscible 50/50 (wt) PS/PMMA blends remain phase separated after being melt-blended with organoclays, the addition of which results in a drastic reduction in the average PS domain sizes (from 1–1.5 μm to ca. 0.3–0.5 μm). Moreover, the majority of clay particles is concentrated in the PMMA phase and in the interfacial region between PS and PMMA. They explained this domain size reduction by the combination of partial compatibilization due to excessive surfactant used in preparing the organoclays and increased viscosity.

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For immiscible 80/20 and 50/50 PS/PP or PS/PP-g-MA blends, Ray et al. [67] clearly indicated that organically modified layered silicate acts at the same time as a nanofiller and also as a compatibilizer. Silicate layers are either intercalated or exfoliated and are located at the interface between the two polymers. Such interfacial activity results in a dramatic decrease in interfacial tension and thus in particle size for both PS/PP and PS/PP-g-MA blends. The compatibilization process is more efficient when PP is grafted with MA that ensures interactions with clay side OH groups. Ray et al. [66] also reported the compatibilization efficiency of the organoclay in an immiscible 60/40 PMMA/PC blend and clearly showed that the key factor for the compatibilization efficiency is the initial interlayer spacing of the organoclay. Si et al. [237] studied the morphologies of PS/PMMA, PC/SAN, and PMMA/EVA blends with and without modified organoclays. In each blend they found a large reduction in domain size and the localization of the clay platelets along the interfaces of the components. Moreover, they demonstrated that compatibility can be enhanced in a wide variety of polymers simply by melt-mixing with modified organoclays. This property is attributed to the large surface area of the organoclays which enables the formation of in situ grafts during melt-mixing. The grafts are then unstable in either one of the phases, and become localized at the interface between the polymer components. Very recently, Bandyopadhyay et al. [289] also showed that the organically modified nanoclay plays the role of a surfactant in the nonpolar rubber/polar rubber blends such as natural rubber/epoxidized natural rubber (NR/ENR) blends. Moreover, the distribution of clay in the rubber phases in rubber/rubber blends is primarily dependent on the nature of the rubbers, their viscosity, and the nature of the clay. In a nonpolar–polar NR/ENR blend, the distribution is predominantly controlled by the clay rubber interactions. Unmodified clay, with a higher polarity, migrates more into the polar rubber phase (ENR), whereas the reverse is true for modified clay.

6.3.6.2

Increase of Matrix Viscosity

The components in polymer blends generally exhibit largely different polarity. In the blends where the matrix phase has strong polarity, organoclay is mainly located in the matrix. In this case, there is a reduction in the ηr , or an increase in the melt viscosity of the clay-containing matrix, which leads to higher stresses to be imposed to the dispersed phases. Li and Shimizu [70] prepared 50/50 (w/w) PPO/PA6 blend nanocomposites by directly melting extrusion of PPO, PA6, and organically modified clay. The dispersed PPO phase size is significantly decreased from 4.2 to about 1.1 μm upon adding a small amount of clay (2%), for which two explanations were proposed by the authors. First, the selective localization of clay in PA6 phase changes the ηr of the PPO and PA6 phases. Therefore, clay has significant effects on the morphology of the blend. Second, the high aspect ratio of the clay platelets may exert significant effects on the phase coalescence during melt-mixing. However, when the organoclay amount is more than 5%, the original matrix-domain morphology is transformed into the co-continuous morphology.

6.3.6.3

Suppression of Coalescence

The resulting changes in the ηr due to the localization of organoclay in the matrix phase are expected to affect the balance between breakup and coalescence of domains. Higher viscosity of organoclay-containing matrix may hinder the coalescence during mixing. On the other hand, the organoclay migrating to the interface between the two phases can also effectively suppress the domain coalescence due to the decrease in the interfacial tension between the two phases. For a 90/10 PE/PBT blend, Hong et al. [75] predicted that the interfacial tension is reduced from 5.76 to 0.14 cN/m when adding 1 wt% of organoclay. This ability of the organoclay to suppress the coalescence effectively reduces the droplet size.

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Figure 6.31 Plots of number-average (Dn ) and volume-to-surface area average (Dvs ) domain diameters of 80/20 (w/w) PA6/EPR blend vs clay amount. Reprinted from [61]. Copyright (2004) with permission from American Chemical Society.

Using melt-blending, Khatua et al. [61] studied the effect of organoclay platelets on morphologies of three blend compositions (80/20, 20/80, and 99.5/0.5 w/w) of PA6/EPR. Figure 6.31 shows the plots of the domain diameters versus the amount of the clay for the 80/20 PA6/EPR blend. It can be seen that the dispersed domain size decreases significantly at lower amounts of the clay, and then slowly but gradually with further increasing the clay amount. This curve is similar to the emulsification curve, which has been reported for an immiscible blend with a compatibilizer, as shown in Figure 6.21. However, when EPR becomes the matrix phase for the 20/80 blend, dispersed PA6 domain does not decrease with increasing the amount of clay up to 2 wt%. Furthermore, for 99.5/0.5 blend dispersed EPR domains does not change with the amount of the clay. The authors concluded that exfoliated clay platelets in the polymer matrix can effectively prevent the coalescence of dispersed domains, resulting in the decreased domains. However, the clay does not improve the interfacial adhesion between PA6 and ERP phases as a compatibilizer does. That is, this domain size decrease does not mean the increase in the compatibility (or miscibility) between two immiscible polymer chains since the organoclay platelets are not located near the interface. The work of Lee et al. [286] on PP/ethylene–octene based elastomer blend nanocomposites containing 30% elastomer revealed two important features for their morphology. Firstly, the aspect ratio of the clay particles decreases as the clay amount increases. Secondly, the elastomer particle size decreases with increasing clay content, which is most likely due to the increase in melt viscosity and the role of clay particles in preventing the coalescence of elastomer particles during mixing. The rate of reduction in both particle aspect ratio and elastomer particle size is higher with initial loading of the clay, and, then, both decrease slowly with further addition of clay. By the intensive mixing, Hong et al. [75, 285] prepared immiscible PBT/PE blend nanocomposites with various compositions of PBT ranging from 1 to 90 wt%. Upon the addition of a small amount of organoclay (1–3 phr), the thin organoclay tactoids with the thickness of the order of 10 nm are located at the PBT/PE interface. This hydrodynamically and thermally stabilizes the blend morphology by the variation of the interfacial tension and coalescence suppression of the droplets. As its content is increased, the additional organoclay selectively locates in the PBT phase due to its affinity with PBT. This results in effective size

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0

10

20

80

90

100

PBT [%]

Figure 6.32 Volume-to-surface-area average diameter (Dvs ) of PBT/PE blends depending on PBT composition. Reprinted from [75]. Copyright (2007) with permission from Springer.

reduction (Figure 6.32) and narrowed size distribution of the dispersed phase. As can be seen in Figure 6.32, the average domain size of the PBT/PE blends increases significantly with PBT content. Upon addition of 4.8% organoclay, the domain size converges to an average diameter of 1.2 μm, regardless of PBT content. In the blend with PBT matrix, the domain size exhibits a significant reduction with the addition of organoclay. Yoo et al. [62] examined the morphologies of LLDPE/PA6 blend nanocomposites with different compositions prepared by melt-mixing. The size of phase-separated domains decreases considerably with increasing organoclay content. The d-spacing of organoclay in the nanocomposites is increased from about 18.6 to over ˚ This effect is highly dependent on PA6 contents because PA6 is more polar and shows higher affinity to 28A. the organoclays compared to LLDPE. Most of the aforementioned results clearly show that the effect of the organoclay on the reduction in the size of dispersed phases is governed by its location in blends. The organoclay-containing component shows an increase in its viscosity, which means the variation in the ηr . If the organoclay is situated in the matrix, the size reduction effect is enhanced as the amount of organoclay is increased. On the contrary, if the amount of organoclay located in the dispersed phases is increased, the deformability and breakup capability of the dispersed phases is significantly reduced and the tendency for coalescence is higher, so there often appears to be an increase in the dispersed phase size [72, 75, 285]. This can be observed in Figure 6.30. However, an enhancement in the breakup of elongated threads of the dispersed phase was reported in the case of the dispersed phase containing non-exfoliated clay [290]. This may seem that the state of clay exfoliation in the dispersed phase determines to a large extent its influence on the final blend morphology. Actually, two of the aforementioned three mechanisms may be evoked simultaneously regarding the phase size reduction when adding the clays. For example, as mentioned previously, the change in the ηr , together with the coalescence suppression effect, affects the determination of the droplet size, depending on the location of the organoclay. In addition, Mehta et al. [287] prepared TPO nanocomposites using melt-mixing. The TPO is composed of about 70 wt% iPP and about 30 wt% EPR. The EPR particle is found to undergo progressive breakup and decreases in size. The number-average diameter of EPR domain decreases by a factor of about 3 as the clay loading increases from 0.6 to 5.6 wt%. The authors ascribed this to one or both

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of the following mechanisms: (1) the melt viscosity increases as the clay loading increases, and this may play a role in the control of the domain size through shear-mixing coalescence and breakup; and (2) the clay preferentially segregates to the EPR particle/matrix interface, and it is suspected that the progressive breakup of the EPR particles is due to the accompanying chemical modifiers of the clay, acting as compatibilizers and reducing the interfacial tension. Furthermore, there is a clear relationship between the clay location and the shape of dispersed phases in the blend nanocomposites. Martins et al. [291] found that for the 60/40 PP/EVA blend where the organoclay is in the dispersed EVA phases, they exhibit a lamellar shape. It is speculated that this shape is due to the presence of a fraction of exfoliated clay particles located at the interface between PP and EVA phases, which leads to the irregular shape of the EVA phases by hindering the rounding promoted by the interfacial tension in the neat PP/EVA blend. Similar lamellar shapes of the dispersed phases were also found in other blends when the organoclays are located in the dispersed phase, such as in the PP/PA6 [72] and HDPE/PA6 [283] blends. Ahn and Paul [64] showed that for PA6/EPR-g-MA blend, the presence of the clay platelets appears to affect the dispersion of the rubber phase resulting in larger and elongated rubber particles. Filippone et al. [284] found that the organoclay appears predominantly confined inside the PA11 phase in LDPE/PA11 blends. This strongly affects the micron-scale arrangement of the polymer phases, causing a drastic refinement of the microstructure for the blends with droplets/matrix or co-continuous morphology and an enhancement of the interfacial adhesion. Recently, several authors [68, 70, 71, 74, 82] have observed that the organomodified clay promotes the formation of co-continuous morphology in polymer blends prepared by melt-mixing. For 50/50 (w/w) PPO/PA6 blend, Li and Shimizu [70] reported that the original matrix-domain morphology is transformed into the co-continuous morphology when the organoclay amount is more than 5%. Later, the same authors [71] also found that for 60/40 ABS/PA6 blend, the typical sea-island morphology for pristine ABS/PA6 blend is changed into a co-continuous morphology on addition of 2% organoclay. When further increasing the organoclay content, the blends keep their co-continuous morphology, but the phase size decreases with increasing clay content. For 60/40 PE/PBT blend, Wu et al. [74] showed that the addition of clay (>2 wt%) changes the phase morphology into a co-continuous one. The sharp increase in the viscosity of the PBT phase due to the percolation of the clay tactoids tied with PBT chains results in a remarkable increase in the ηr between the PBT and PE phases, and so promotes the phase inversion. This may be a possible mechanism in the formation of the co-continuous morphology in the blend. Ray et al. [68] prepared 80/20 PP/PBSA and PP-g-MA/PBSA blends upon addition of 5 wt% organoclay. For the former blend, the organoclay changes the morphology from highly phase-separated to a typical co-continuous structure. This is attributed to the selective localization of intercalated silicate layers in the PBSA phase, which increases the viscosity of the PBSA phase. On the other hand, the addition of organoclay into the latter blend reveals the efficient mixing of the polymer matrices. This is due to the common intercalation of both polymer chains into the same silicate galleries. Using 75/25 HDPE/PA6 blend, the research by Filippone et al. [82] showed that a shigh degree of cocontinuous morphology can be formed at low contents of dispersed phases when adding the organoclay. They found that the extrusion process promotes the formation of highly elongated and separated organoclay-rich PA6 domains interspersed with the host HDPE. The morphology of the extruded sample, however, evolves to a co-continuous morphology when the extrudate is melted again. The degree of co-continuity is increased from 22.8% for the pristine HDPE/PA6 blend to one larger than 87.3% upon addition of 4.8 wt% clay, which clearly indicates that the PA6 in the blend is characterized by a very high degree of continuity. Besides organoclays, some other nanoparticles also can influence the morphology of immiscible polymer blends. Li et al. [288] prepared PP/PET/TiO2 nanocomposites by compounding a PP/TiO2 nanocomposite premix with PET in absence and presence (up to 6 vol%) of PP-g-MA. The TiO2 nanoparticles are mainly located at the PP/PET interface and to a lesser extent in the dispersed PET droplets without PP-g-MA. The TiO2 nanoparticles are dispersed in the PP matrix as well as at the interface and are exclusively located

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in the PP upon addition of 3 and 6 vol% of PP-g-MA, respectively. The incorporated TiO2 nanoparticles result in finer morphology of the PP/PET blend irrespective of their locations. Locating at the interface, the TiO2 nanoparticles exhibit a remarkable compatibilizing effect, by decreasing the free energy of mixing and preventing the coalescence of PET droplets. Whereas preferentially located in the PP matrix, the TiO2 nanoparticles decrease the viscosity ratio which facilitates the breakup of PET droplets. Zhang et al. [292] studied the morphology and electrical properties of PA6/PP/MWNTs composites with varied compositions. The MWNTs were found to be located preferentially in the PA6 phase, and a small amount of MWNTs bridges the PA6 and PP phases, whatever the PA6 is continuous or dispersed phase. The incorporation of MWNTs transforms the dispersed PA6 phase from spherical to elongated or irregular shape, which exhibits a higher electrical conductivity. Very recently, Baudouin et al. [293] demonstrated the possibility to confine a large fraction of unpurified and unfunctionalized MWNTs at the interface of an immiscible PA/EA blend, taking advantage of irreversible polymer adsorption on MWNTs and their subsequent migration. Moreover, the final morphologies observed are stable. Contrary results, that is, the nanoclay increases the size of dispersed phase, were reported. Via melt processing, Yoo et al. [294] prepared 80/20 aPA/elastomer (EOR or EOR-g-MA) blends. They found that most of the organoclay is well exfoliated in the aPA phase, but some of the organoclay migrates to the interface and tends to envelop the EOR phase. The average elastomer particle sizes in aPA/EOR blends decrease significantly upon adding organoclay. It seems that the addition of organoclay to such blends can be an effective way of tailoring elastomer particle size and stiffness enhancement; whereas for a-PA/EOR-g-MA blends, the addition of organoclay slightly increases the elastomer particle sizes due to the fact that the elastomer particle sizes in such blends are already very small. The study by Gallego et al. [295] showed that for PA6/mEPDM/EPDM-g-MA blends, the mEPDM particle sizes increase with increasing amount of the nanoclay (2, 3, and 4 wt%).

6.4

Processing-Relevant Factors Affecting the Morphology

The processing-relevant factors that would affect the morphologies in multiphase polymer systems mainly include flow field types, processing parameters (such as shear/elongational rate or shear/elongational stress, mixing temperature, and mixing time), and the mixing sequence. As mentioned in Section 6.2.1, the most significant morphology development of polymer blend occurs at the very early stage of mixing and further mixing often has no or little influence on its morphology. This, however, does not necessarily mean that the final morphology is determined mainly in the solid–melt transition stage. The final morphology may be largely controlled by the geometries or the flow fields in the metering section of the extruder and in the extrusion die, as well as processing parameters [28, 276, 296, 297]. 6.4.1 6.4.1.1

Flow Field Types Elongational vs Shear Flows

The flow in polymer blending equipment is a mixture of both shear and elongational flows. Concerning droplet breakup and dispersion in polymer blends, elongational flows are much more effective than simple shear flows over a much wider range of viscosity ratios [121, 134, 136]. In elongational flow, the most effective region of ηr for dispersing particles is between about 10 and 40, and the Cacrit is lower than that in simple shear flow [134, 298]. Heindl et al. [298] believed that it is possible to incorporate the true hydrodynamic stress for calculating the capillary number during uniaxial elongational flow by considering the transient elongational viscosity of the matrix (η+ E,m ). By replacing the shear viscosity of the matrix and shear rate in Eq. (6.2) by the

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η+ ε ), respectively, they defined a capillary number for blends undergoing elongational E,m and the strain rate (˙ flow, CaE , as follows: Ca E =

η+ ˙ E,m ε 2σ/D

(6.29)

Thus, the experimentally found deformation behaviour of the dispersed phase of blends can be related to the CaE at different ε˙ and temperatures. Using a converging channel is the simplest way to introduce uniaxial elongational flow during polymer processing. Bourry et al. [118], Khayat et al. [119], and Godbille and Picot [120] experimentally and numerically studied the influence of elongation and shear on drop deformation and breakup in a convergingdiverging channel. During the convergent flow, the Newtonian drop is initially reluctant to deform, but then its deformation is rapid. By contrast, the viscoelastic drop initially readily deforms, but then its deformation slows down. For the viscoelastic drop, its initial diameter (D0 ) has a dramatic effect on its deformation process and breakup mechanisms. On the centerline of the channel, where the flow is purely elongational, the drop deformation is independent of the D0 for a lower value of ηr . As the ηr increases, the deformation increases with increase of D0 and with decrease of ηr . Smaller drops deform symmetrically, but do not break up. Larger drops show a horseshoe deformation after passing through the narrowest zone of the channel and then break up by two distinct mechanisms: (a) drop pinching at the parallel channel entrance, creating a tail which breaks up quickly to form 20–100 μm droplets, and (b) subsequent tip streaming at the downstream of the pouch-like fluid disc forming in the parallel channel, producing 5–15 μm droplets. Furthermore, for the larger drops and higher values of ηr there is a loss of shape recovery (symmetry). The magnitude of the drop-to-matrix deformation is in a qualitative agreement with small perturbation theory. Small drops located away from the channel centerline undergo complex deformation caused by shear and elongational stresses. Also using a converging channel, Meller et al. [121] studied the dispersive mixing of blends with high values of ηr (8–450). They found that the dispersive mixing efficiency is proportional to the maximum elongational stress within the converging section, and to the length of the region where the critical conditions for elastic fracture of droplet material are met. Dispersive mixing mechanism in blends with high values of ηr is mainly dictated by elastic fracture of the droplet, and hence is applicable over a wide range of polymer blending processes. A four-roll mill can be used to generate a planar elongational flow. Using this device, Ha and Leal [123] investigated the flow-induced stretching of Newtonian droplets in a Newtonian or a viscoelastic (Boger) fluid at high capillary numbers. If the stretch ratio is large enough, the droplets undergo an end-pinching process and break into two or more parts with a small satellite in between upon cessation of flow. Moreover, the elasticity of the suspending fluid seems to stabilize the elongated droplet against breakup. The observation by Handge [125] showed that in equibiaxial elongation, the initially spherical PS drop embedded in a PMMA matrix is deformed into a very flat circular disc. In the initial stage of relaxation, holes are formed near the rim of the drop and increased in number and in size with time. The formation of holes resembles that during the initial stage of blending of two immiscible polymers in the melt [30, 85, 86, 89, 101, 299, 300]. At the later stage of relaxation, the original PS drop attains a fibrillar shape. Interfacial tension driven breakup of an equibiaxially elongated drop can lead to a much larger number of small droplets than the breakup of a uniaxially stretched drop. Also using a four-roll mill, Tretheway and Leal [124] studied the deformation and relaxation of a Newtonian drop (PDMS) suspended in a purely elastic PIB/PB fluid. Their results suggest that a nonlinear coupling exists among the steady-state drop shape, the local disturbance flow, and the polymer configuration in the vicinity of the drop. As a consequence of this coupling, the drops become more deformed, with ends that are generally

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more pointed. The inverse problem of an elastic PIB/PB drop in a Newtonian (PDMS) fluid in planar elongational flows was studied by Hsu and Leal [126], who showed that the viscoelastic effects alter the steady drop shape from being ellipsoidal to drop shapes with more blunt ends. It has been reported that a perfect microfibrillar morphology is developed when an effective elongational flow is applied to a blend of two polymers with a suitable value of ηr [301]. Several researchers have showed that exerting a sufficient elongational flow can create fibrillar morphology for blends with suitable values of ηr [169]. By quenching the samples, Gramespacher and Meissner [302] showed that the originally spherical PS droplets in the continuous PMMA are stretched into ellipsoids and finally into long fibrils during uniaxial elongation. Experiments by Lacroix et al. [301] showed that the PP/EVA/EMA blends, after extrusion through a hyperbolic-shaped die exhibit fibrillar morphology, due to the elongational component of the flow field. By using several models, it is possible to qualitatively predict the droplet-fibril transition. In general, elongational flows are more effective for this transition than shear flows. Through die flow experiments, Oosterlinck et al. [303] demonstrated that the PS droplets deform into fibrils by the uniaxial elongational flow in the entrance region of the capillary. Whether fibrillar morphologies break up in the capillary depends on their shear history and hence on their radial position. Oosterlinck et al. [304] found that during uniaxial elongational flow the droplet deformation causes an extra stress related to the interfacial tension. This extra stress can be calculated from the droplet deformation. So, in principle it is possible to use rheological measurements to probe the underlying morphology development during uniaxial elongational flow. Star´y and M¨unstedt [305] showed that during uniaxial deformation, the dispersed LLDPE droplets stretch and form highly elongated fibrils in agreement with a modified capillary number model (Eq. (6.29)). Moreover, the droplets are deformed less than the surrounding matrix. This is probably caused by a slip at the interface. The flow field type in polymer blending equipment may also exert an obvious influence on the microstructure of polymer nanocomposites. Huang and co-workers showed that high shear intensity [306], especially the distributive mixing [307, 308] provided by the TES facilitates to the size decrease and dispersion of the nanoCaCO3 in PP matrix. Some investigators, e.g. Gupta et al. [309], observed that upon uniaxial elongational flow the microstructure changes from a mainly exfoliated state to some sort of re-aggregation of the silicate platelets, and hence the re-formation of tactoids in polymer/clay nanocomposites. This may be due to the stretching rate components perpendicular to the velocity, whereas Utracki et al. [310, 311] showed that the elongational flow does not cause re-aggregation of the clay particles reported for shear. Preliminary investigation on LLDPE/clay nanocomposite drawn fibers by La Mantia et al. [312] suggested that elongational flow can induce exfoliation of intercalated tactoids and some more intercalation of the same tactoids. Moreover, broken clay particles are oriented along the flow direction. From this point of view, polymer nanocomposites can be considered as polymer blends. Further, La Mantia et al. [313] demonstrated that the sizes of the clay particles are more and more reduced with increasing applied elongational stress. Garofalo et al. [314], aiming at investigating the effect of elongational flow on the nanoscale arrangement of the silicate inside PA-based nanocomposites, demonstrated that isothermal and nonisothermal elongational flow can modify the nanomorphology of the nanocomposites enhancing exfoliation degree and promoting orientation along stretching direction. Moreover, applied stretching induces clay–clay electrostatic interactions, which can translate into a three-dimensional silicate arrangement and/or aggregation of the filler. The entity of stretching-induced structural modifications is highly dependent on the initial nanomorphology, polymer–clay affinity, and applied stretching. Handge and P¨otschke [315], focusing on the elongational properties and the morphology of an electrically conductive PC/MWNTs composite, also showed that isolated carbon nanotubes are aligned in elongation. Moreover, the arrangement of carbon nanotubes results in a yield stress and prohibits large extensions of the macromolecules during elongation. Utracki and co-workers [118, 316, 317] developed an extensional flow mixer, EFM, a device in which the fluid flows through a series of convergent/divergent regions of increasing intensity. It was observed for most cases that a SSE equipped with an EFM (SSE + EFM) outperforms a TSE in generating finer dispersion,

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improving the blends’ toughness, or facilitating dissolution of very high molecular weight fractions. Moreover, better dispersion was found compounding the polymer/clay nanocomposites in (SSE + EFM) than in TSE with or without EFM [310, 311]. 6.4.1.2

Complex Flows

In real polymer mixing processes, the flow field is not steady shear or steady elongational flow, but is always a mixed or complex flow. By analyzing the time scales of various breakage mechanisms studied in simple shear flow and comparing them to the appropriate time scales in more complex flows, Bigio et al. [137] predicted the relevant importance of a given mechanism in determining the daughter drop size distribution. They carried out this type of time scale analysis for flow in an extruder and showed that the mechanism can be important when axial flow contributes to the strength of the shear forces acting to disrupt the drop. Using more concentrated model (PIB/PDMS) blends, Testa et al. [318] visualized morphology under complex flow fields combining shear and elongational flow, which was clearly shown to be strongly affected by the complex nature of the flow. The flow-induced morphology development can be predicted (at least semi-quantitatively) by assuming that the effects of elongational and shear components are additive. The criterion for droplet breakup derived from these experiments is that breakup occurs when either the local shear or stretching rate at a certain axial position exceeds the critical value for breakup as obtained from steady homogeneous flow data. For strong drop deformation, generating slender filaments, a sudden decrease in shear rate leads to breakup. Using an eccentric cylinder device, Boonen et al. [132] investigated the deformation and orientation of single Newtonian droplets dispersed in a Newtonian matrix undergoing sub-critical complex flows. They compared experimental results with predictions obtained using the transient form of the phenomenological model of Maffettone and Minale [319] and found good agreement between both. So, their study provides a quantitative assessment of drop shape predictions in complex flows. For PS/clay nanocomposites, Nassar et al. [320] found that a combined shear and elongational flow field results in a more effective dispersion of nanoclay in matrix than a pure shear or elongational flow. Moreover, a single droplet in a complex flow field experiences continuous, transient rates of shear and elongation as it moves along its trajectory through the flow field [150]. The effect of transient flow on drop deformation is inherently more difficult to understand and categorize. Numerical and experimental investigations of transient effects on the deformation and breakup of drops in simple flow fields include those of Delaby et al. [321, 322], Toose et al. [323], and Hooper et al. [324]. 6.4.2

Chaotic Mixing

In most polymer processing equipment, there is only one possible mode of flows–laminar flow. While only laminar flow is available for mixing, chaotic flow is known to provide the best possible mixing. Beginning in the early 1990s, it was demonstrated that polymer blend morphology can be developed progressively by chaotic mixing. Chaotic mixing has been used to more controllably organize melt components and solid additives into a variety of fine-scale morphologies for improving properties or imparting functionality of polymers during recent years [325]. 6.4.2.1

Introduction of Chaotic Mixing Characteristics

A flow t (•) in a region R displays chaos when it satisfies one of the following conditions [326]: (1) there is an invariant set R [i.e., t (R) = R] and the flow is sensitive to initial conditions on R; (2) the flow has transverse homoclinic and/or heteroclinic points; (3) the flow produces horseshoe maps. Among the three definitions, horseshoes offer the only direct experimental verification of chaos.

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Initial minor component body

Characteristic 1: Stretching and folding

Characteristic 2: Sensitivity to initial locations

Initial cluster of particles in masterbatch

Figure 6.33 Schematic of two defining and related characteristics in chaotic flow. Reprinted from [325]. Copyright (2006) with permission from John Wiley & Sons.

Excellent mixing ability of chaotic flow is attributed to its two defining and related characteristics [130, 325], which are shown in Figure 6.33. In Characteristic 1, the region of space enclosing an initial minor component body becomes stretched and folded in response to the periodic motions of bounding surfaces. In Characteristic 2, the positions of individual parts such as particles in a masterbatch of an initial minor component body diverge exponentially with time. As a result, a chaotic flow can generate a uniform spatial distribution of the components and fine-scale microstructures. Ottino et al. presented a systematic approach to the modeling of distributive mixing by combining the kinematical foundations of fliud mechanics with chaotic dynamics. Of special interest is their numerical and experimental work on chaotic flows mainly in two-dimensional geometries [127]. Some pioneers have applied tools of dynamics to study the chaotic mixing in polymer processing. Kim and Kwon [327] carried out the numerical simulations of chaotic flow in Chaos Screw in terms of particle trajectories, Poincar´e sections, and mixing patterns, along with the invariant manifolds. Cheng and Manas-Zloczower [328] numerically simulated the distributive mixing in single and counter-rotating TSEs with emphasis on the chaotic features of flow in them. In their work, Lyapunov exponent is used to quantify the divergence of initial conditions. Positive values of this exponent indicate an exponential growth in the length and area stretch, that is, a certain degree of chaos for the flow field. For a given motion, the Lyapunov exponent is (λL )defined by: λL =

lim

t→∞

x(0)→0

1 | x(t)| ln | x(0) | t

(6.30)

| x(t)| where | x(0)| represents the specific rate of stretch. Their results showed that the SSE exhibits a certain degree of chaotic flow in terms of calculated Poincar´e sections and Lyapunov exponents. Moreover, a comparison

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of the Lyapunov exponents for different extruders or different operating conditions correlates well with the degree of distributive mixing achieved in the system as expressed by average length stretch values. 6.4.2.2

Various Apparatuses for Generating Chaotic Mixing

To investigate the dynamics of stretching, folding, and breakup of droplet in a chaotic flow field, an eccentric cylinder device was developed [127, 128, 130, 329]. In this device, a time-periodic two-dimensional chaotic flow field is created by altering co-rotation and counter-rotation of the cylinders. Both batch and continuous mixers have been developed to generate the chaotic flows in polymer processing. In the batch chaotic mixer developed by Zumbrunnen and co-workers [330–332], two- or three-dimensional chaotic mixing is induced by alternately and repetitively rotating the inner and outer cylinders, whereas the cross-sectional shape of the mixing chamber in the batch chaotic mixer developed by Jana and co-workers [333, 334] is similar to that of TSEs. This design can provide larger and more uniform stretching of the material interfaces by adjusting the nature of time-periodic waveforms and degree of chaotic mixing. For industrial applications, continuous chaotic mixers should be used. A continuous chaotic advection blender (CCAB), as shown in Figure 6.34, was developed by Zumbrunnen and co-workers [335]. It consists primarily of two stir rods of circular cross section with tapered ends enclosed in an oval barrel. Polymer melts A (minor component) and B (matrix) from the SSE and metering pumps enter the CCAB via a cylindrical melt distribution block, which contain nine smaller ports for polymer A and one central larger port for polymer B. The rods are rotated by variable speed motors. For continuous screw extrusion process, Huang and Peng developed a convective screw, called the HP Screw, which was shown to induce dispersive melting [336] and chaotic mixing [307]. Kim and Kwon [337]

Variable Speed Motor & Reduction Gear

Polymer A Progressive Morphology Tapered Circular Stir Development Rods in Oval Barrel

Extrusions

Melt Distribution Block Metering Pump Polymer B Process Control Computer

Figure 6.34 Schematic of continuous chaotic advection blender (CCAB). Reprinted from [335]. Copyright (2005) with permission from Elsevier.

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enhanced the mixing performance of the SSE by inserting periodic barriers in the screw channel and named it the Chaos Screw (CS). The studies showed that the domain-averaged Lyapunov exponent in CS is positive, indicating the flow in CS is chaotic. The main features of both HP and CS Screws are the simple geometry and the facility in machining, installation and maintenance. Moreover, chaotic mixing also exists in static mixers due to the splitting and reorientation of fluid flow [338]. 6.4.2.3

Deformation and Breakup of Droplets in Chaotic Mixing

Using the aforementioned eccentric cylinder device, Ottino and co-workers investigated the dynamics of stretching, folding, and breakup of droplet in a two-dimensional chaotic flow field [127, 128, 329]. Tjahjadi and Ottino [128] showed that droplets stretch into long and axisymmetric filaments, which stretch, extend, fold upon themselves nonuniformly over time and eventually break up into many smaller droplets; then chaotic stirring disperses the fragments throughout the flow. Droplet breakup near folds is associated with a change of sign in stretching rate; this breakup mode results in the formation of rather large droplets. The dominant breakup mechanism, however, is capillary wave instabilities in highly stretched filaments. Other breakup modes, such as necking and end-pinching occur as well. The average droplet size decreases as the ηr increases. This is attributed primarily to the fact that higher droplet viscosity delays the onset of capillary breakup. Thus, in low ηr systems (0.01–1), droplets stretch relatively little before breaking up, resulting in the formation of larger droplets that may or may not break again. In higher ηr systems (1–2.8), droplets stretch substantially before breakup, producing very small fragments that rarely break again. This leads to a more nonuniform equilibrium droplet size distribution. It was also observed that the Cacrit for a thread radius to produce breakup is often lower than that for the original droplet. The implication of this is that the transient breakup mechanism is more effective than repeated droplet breakup at the Cacrit , if the primary aim is to produce the finest system morphology and the smallest droplet sizes. Muzzio et al. [329] found that deformation and breakup of immiscible fluids in a chaotic flow is governed by self-similar distributions of stretching histories and stretching rates and produces populations of droplets of widely distributed sizes. Droplet size distributions upon scaling fall into two self-similar families of curves, depending upon the ηr . Each family exhibits a different distribution shape, presumably due to changes in the breakup mechanism. Florek and Tucker [130] examined the stretching behavior of droplets with high viscosity and zero interfacial tension in two-dimensional, time-periodic chaotic flow between eccentric cylinders. High-viscosity droplets stretch slowly in the first few periods, but progressively become elongated and oriented in directions favorable to subsequent stretching. Eventually the droplet elongations and orientations follow the stretching of the underlying flow. Consequently, at long times, droplet stretching statistics follow the universal features of passive fluid elements in a chaotic flow: the geometric mean stretch grows exponentially at the rate of the Lyapunov exponent, while the log of the principal stretch ratio, when scaled by its global mean and standard deviation, exhibits an invariant global probability distribution and an invariant spatial distribution. These results demonstrate that chaotic flows are highly effective at stretching ‘hard to stretch’ morphologies. The stretching ability of a globally chaotic flow can be characterized by a modest amount of information: the Lyapunov exponent, the growth rate of the variance of the stretching exponent, and the asymptotic global probability distribution and spatial distribution of normalized stretch. The main effect of increasing the ηr is to slow the initial stretching rate and to increase the time required to reach exponential stretching and asymptotic statistics. This work was extended by Pham and Tucker [131] to include interfacial tension and its effect on droplet shape. A numerical simulation was developed to study the time-dependent shapes of droplets in chaotic mixing as a function of interfacial tension and ηr by the authors. A Lagrangian particle method was used to follow the morphology, which is modeled as three-dimensional ellipsoidal droplets, ignoring breakup and coalescence. Droplet stretching behavior depends on interfacial tension and viscous stresses. In the regime of the global capillary number (Cag ) where the major axes of the droplets stretch exponentially,

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the behavior of the minor axes is sensitive to the interfacial tension. For lower Cag in this regime, interfacial tension quickly forms the droplet cross-section into a circle, and the droplet assumes a thread-like shape. For higher Cag , initially spherical droplets are first stretched into ribbons or sheets, where the minor axes in the flow plane contract, while the axes in the neutral direction change very little. The higher the Cag , the longer the ribbon shapes exist and the thinner they become. However, the ribbons eventually relax into threads. If the goal is to produce a sheet-like morphology, the best strategy is to use a chaotic flow with high speed and a large initial size of the dispersed phase. This promotes the largest possible Cag . 6.4.2.4

Morphology Development of Multiphase Polymer Systems in Chaotic Mixing

Both the batch and continuous chaotic mixers described in Section 6.4.2.2 have been employed to prepare polymeric blends and nanocomposites. Various structures, such as lamellas, fibrils, droplets, and their combinations were developed in the blending of PS with LDPE [299, 331, 332, 339], PP with EPDM [332], LDPE [335] or PA6 [84, 333, 334], and PA or LDPE with EVOH [340, 341]. Zumbrunnen et al. systematically investigated the morphologies induced by chaotic mixing and their relationship with resulting mechanical properties for PS/LDPE, PP/EPDM, and PP/LDPE blends. When using the batch chaotic mixers, they found that for N (numbers of chaotic mixing periods) <20, lamellar structures exhibiting high impact toughnesses and long fracture times are produced. When 20 < N < 30, fibrillar structures providing greater improvements in impact properties arise [331]. Using a continuous chaotic mixer, multilayer film morphologies were produced with layer thicknesses below 200 nm in PS/LDPE blends. Long and unbroken LDPE films develop quickly and become distributed. The PS major component is similarly converted into film form as both LDPE and PS are stretched and folded about one another. The LDPE layers also become interconnected to result in an interpenetrating blend (IPB) [332]. The achievable reduction in layer thickness and the corresponding increase in layer number by chaotic mixing depend primarily on interfacial tension and ηr . Layer thicknesses of less than 200 nm were also produced in PP/EPDM (80/20 vol) blend where interfacial tension is smaller [332]. PS/LDPE blends were also selected to investigate the morphology development mechanism via the batch chaotic mixer [299]. The transition from a sheet morphology to an IPB morphology is illustrated in Figure 6.35. Various types of hole formation and growth play a central role in this morphology development. Holes form at locations in layers where thickness is smallest and enlarge preferentially in the direction of least thickness. Hole formation occurs interactively in the multilayer blends such that phase continuity occurs among previously separated layers. Layered morphologies yield interpenetrating blends over a broad compositional range. Hole enlargement subsequently causes layers to fragment into abundant platelets to give blends with single phase continuity. For LDPE/EVOH blends, Kwon and Zumbrunnen [341] found that optimal barrier properties are obtained in a novel single phase continuous and mechanically interlocked morphology that is an outcome of stretching and folding in chaotic mixing. Chaotic mixing is effective in developing a desirable morphology consisting of distributed and numerous thin platelets with submicron layer thicknesses and other thin layer fragments. These shapes impart a high tortuosity and increase the crystallinity, so that extruded films have barrier properties close to those of films containing unbroken thicker layers. Moreover, higher barrier properties correlate with smaller layer thicknesses resulted from chaotic mixing. Jana et al. [333, 334] blended PP with PA6 using the batch chaotic mixer. It was demonstrated that selfsimilar mixing structures, a novel feature of chaotic mixing, are generated as a precursor to an array of mixing microstructures, such as nested layers, elongated fibrils, droplets, and their combinations. A very large fraction (>90%) of droplets generated fall below the equilibrium size and are much smaller than those developed in TSEs. The droplet size distribution does not follow proper scaling probably due to a lack of scalability of interfacial tension driven the breakup of the lamellas.

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Figure 6.35 Interactive hole formation in blends with multilayer morphology and emergence of an IPB: (a) incipient hole formation in region of smallest layer thichness; (b) hole enlargement and adjacent layer thinning; (c) incipient hole formation in layers of other component; (d) multiple holes leading to phase continuity. Reprinted from [299]. Copyright (2002) with permission from Elsevier.

Chaotic mixing offers a means to indirectly manipulate additives with nanoscale dimensions to form a variety of structures. An example is shown in Figure 6.36, where single wall nanotubes orient in a PP matrix. Moreover, the investigation of the effect of layered silicate clay on the morphology development of PP/PA6 blend in chaotic mixing showed that the clays help to produce droplets with much smaller size and narrower size distribution due to their direct influence on the breakup of PP domains [290]. Using a batch chaotic mixer, an electrically conductive PP/PA6 blend was obtained by adding conductive carbon black (CB) particles [342]. The migration of CB affects the conductivity of blend. PP-phase containing CB particles deformed into lamellar and fibrillar morphologies produce continuous networks in the flow direction, and provide conductivity by double percolation. Double percolation networks persist even after fibrils break into droplets aided by migration of CB particles from the bulk of PP droplets to the interface. On continued mixing, the blend eventually turns into insulator as CB particles migrate from the PP/PA6 interfaces to PA6 phase. Electrically conductive PMMA/carbon nanofibers composites were prepared in a low-shear chaotic mixer [343]. Conductive networks are formed early due to orientation of the fibers pulled out of the agglomerates. Electrical conductivity exhibits great sensitivity to mixing time around percolation threshold and remains almost unchanged with prolonged chaotic mixing above the percolation threshold. Huang et al. prepared the PP/nanoclay [344, 345] and PP/nano-CaCO3 [306] nanocomposites by feeding the compounds prepared using a TSE to a SSE with a tunable flow field. The results showed that the chaotic

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Figure 6.36 Single wall nanotubes distributed and oriented by chaotic advection in PP. Reprinted from [300]. Copyright (2002) with permission from American Chemical Society.

mixing provided by the latter favors a greater extent of exfoliation of the nanoclay and the size reduction and dispersion of the nano-CaCO3 . A static mixer attached to the exit of the SSE with a chaotic mixing can further promote the exfoliation of the nanoclay [345], due to the fact that the static mixer exhibits a chaotic mixing feature, as noted previously. Using the continuous extrusion method, Huang and co-workers compared the effects of chaotic and shear mixings on the morphology development for polymer blends [20, 84] and their nanocomposites [84]. It was demonstrated that for the PP/PA6 blend, the PA6 domains are transformed from large particles to short striations and to small droplets finally in shear mixing, whereas morphology of the dispersed phase develops through a wide route involving transitions from lamellar layers to thick striations and to a partial continuous structure finally in chaotic mixing. For (PP/clay)/PA6 blend nanocomposite, the PA6 domains are deformed gradually from droplets to irregular fibrils in shear mixing and from short layers to thin fibrils in chaotic mixing. The PA6 fibrils formed finally in the latter are much thinner and more uniform than those in the former. Moreover, the clay platelets dispersed initially in the PP phase migrate into PA6 phase finally in both shear and chaotic mixing and the exfoliation of clay platelets in PA6 phase is obviously improved in the latter. Figure 6.37 schematically describes the morphology development of this blend nanocomposite in shear and chaotic mixings. 6.4.3

Mixing Sequence

As noted previously, the fact that most significant morphology changes are accomplished during the initial stage of blending of polymer components does not necessarily mean that the final morphology is determined mainly in the solid–melt transition stage. The results by Huneault et al. [208] showed that for polyolefin blends with high values of ηr , successive extrusion passes significantly decrease the number of gels, indicating that the dispersion is not in its final state after a first or even a second extrusion pass. The investigation by Bourry and Favis [346] demonstrated that the final blend morphology depends mainly on the later mixing conditions after both components are molten, even at high contents of the dispersed phase. In their study, three HDPE/PS blend compositions (90/10, 70/30, and 30/70) were mixed in a TSE via three procedures: (1) both HDPE and PS pellets were simultaneously added in the hopper of the TSE; (2) the HDPE pellet was fed at the hopper and the PS melt was injected at the midway of the TSE using a SSE; (3) the PS melt was added even farther

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Figure 6.37 Schematic for morphology development of (PP/clay)/PA6 blend nanocomposite in shear and chaotic mixing. Reprinted from [84]. Copyright (2009) with permission from Taylor & Francis.

downstream with a residence time in the TSE of approximately 25 s. It was shown that the mixing procedure does not significantly affect the size or shape of the dispersed phase. Even when the mixing time is as short as 25 s, the blending of PS and HDPE melts yields the same morphology as mixing pellets over the entire length of the TSE. Furthermore, this morphology is obtained so rapidly in the melt during mixing that it is the last flow field experienced by the blend melt that controls its final morphology. For compatibilized polymer blends, Lin et al. [32] found that the blending sequence determines the initial interfacial distribution of compatibilizer and thus, influences their final morphology. The mixing sequence may have an obvious effect on morphology and properties of ternary blends. Using a TSE, Ha et al. [347] prepared iPP/mPE/HDPE ternary blends via two blending sequences: the simultaneous blending of the three components (method I) and the preblending of mPE and HDPE followed by blending with PP (method II). The results showed that the HDPE particles are encapsulated by mPE in the PP matrix regardless of the blending sequence. Method II results in a much finer morphology than method I. In method I, mPE and HDPE are first separately dispersed in PP, and this is followed by the encapsulation of HDPE by mPE in a core (HDPE)–shell (mPE) morphology. This gives rise to an additional increase in the domain size to the one originally determined by the viscosity ratio of mPE to PP. In method II, however, the domain size is determined by the viscosity ratio of the premixed mPE/HDPE binary blend to PP, which is close to or rather smaller than that of mPE to PP. For ternary PP/EPDM/HDPE blends, Tchomakov et al. [36] found that a two-step mixing procedure where HDPE and EPDM are mixed together before their incorporation into the PP matrix results in finely dispersed composite droplets, in which the HDPE is localized within the EPDM phase. This produces an additional 50% increase in impact resistance when compared to the standard one-step mixing. Huang et al. [348] prepared ternary PA/elastomer blends via three mixing sequences: simultaneous mixing, premixing of the two rubber phases, and master batch preparation. They reported that the mixing sequence seems to cause a negligible difference in the average size of the particles or their polydispersity for blends with a unimodal particle size distribution; however, for blends having a bimodal particle size distribution, the mixing sequence seems to affect the rubber particle size a little more but still not significantly. Various researchers have studied the effect of mixing sequence on microstructure and properties of polymer nanocomposites during recent years. Generally, the mixing sequence also significantly influences the microstructure of the prepared hybrids. Contrary results, however, have been reported. For PA/maleated elastomer/organoclay ternary nanocomposites, such as PA66/SEBS-g-MA/organoclay (80/15/5) [53] and

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PA6/mEPDM/EPDM-g-MA/organoclay (75/10/10/5 or 75/10/5/10) [349] nanocomposites, the mixing sequence shows a high influence on the dispersion of the organoclay and the elastomer. The presence of the exfoliated organoclay platelets in the PA matrix or in the interface between both phases rather than in the dispersed elastomer phase is the best way to achieve equilibrium between stiffness and toughness, which is obtained by premixing the PA6 with the organoclay initially and later blending with the elastomeric component. A different phenomenon, however, has been reported. The addition of organoclay into the PA6/EPR-g-MA blends results in larger and more elongated rubber particles when blending PA6 and organoclay initially and later mixing with EPR-g-MA [64]. Zhang et al. [65] also found that when blending a premixed PA6/organoclay nanocomposite with EPDM-g-MA, the addition of organoclay reduces the viscosity of the matrix and weakens the interfacial interaction between PA6 and EPDM-g-MA, resulting in larger dispersed phase domains. For PA6/mSEBS/organoclay nanocomposites, Gonz´alez et al. [350] found that mixing PA6 with mSEBS and subsequent incorporation of the organoclay results in a decrease in the rubber particle size indicating improved compatibilization and a significant increase in the impact strength. This is attributed to the substitution of the surfactant/mSEBS interaction by the PA6/mSEBS reaction occurring in this blending sequence. Filippi et al. [351] showed that adding the organoclay to the premixed molten LDPE/PA6/SEBS-g-MA blend and using an SEBS-g-MA/organoclay ratio of 0.5 or, preferably, 1.5 allows to obtain an excellent dispersion of the PA6 droplets within the LDPE matrix. For the PP/EVA/organoclay (60/40/5) nanocomposites, Martins et al. [291] found that the blending sequence effect on the dispersion and location of the organoclay in the blends is dependent on the kind of organoclay used. Moreover, a drastic increase in the impact strength and a considerable alteration of the shape of the dispersed EVA phase are observed when the organoclay is located in EVA phase due to the addition of organoclay to the blends. Some studies, such as [57, 286, 352], found that better clay exfoliation and dispersion is obtained by employing an approach to melt-blending organoclay with compatibilizer followed by dilution with matrix (e.g. PP). Such an approach affords additional time for the compatibilizer to penetrate the clay galleries and exfoliate the silicate platelets prior to introducing the matrix resin. Li et al. [353] and Wu et al. [354] prepared PBT/EVA-g-MA/organoclay and PBT/epoxy (as a compatibilizer)/organoclay ternary nanocomposites, respectively. They revealed that the blending sequence in which mixing the EVA-g-MA or epoxy and organoclay first and then mixing it with PBT matrix is the best way to obtain a good dispersion of clay in the matrix. This, along with fine ‘sea-island’ morphology of PBT/EVA-g-MAH blend, results in the best tensile and impact strengths for the former nanocomposite. For the latter nanocomposite, a percolated tactoids network with the highest density and intensity is formed. For co-PP/organoclay nanocomposites, however, Akbari and Bagheri [355] found that premixing of organoclay and compatibilizer does not provide better dispersion of nanoclay in the matrix. Prakashan et al. [356] prepared PP/PDMS elastomer/nano-SiO2 ternary composites with three different mixing sequences. A one-step mixing sequence of all three components results in a predominantly separated dispersion of the PDMS and nano-SiO2 phases in the PP matrix, with some nano-SiO2 particles forming a fine band surrounding the PDMS domains, whereas a few of them got encapsulated inside the PDMS domains. A two-step mixing sequence, where the nano-SiO2 is added in the second step to the PP/PDMS blend, produces the predominant encapsulation of nano-SiO2 by PDMS. Another two-step mixing sequence, in which the nanoSiO2 is added in the second step to the PP/nano-SiO2 binary system, results in separately dispersed PDMS and nano-SiO2 phases in the PP matrix with a loose network of nano-SiO2 particles surrounding the PDMS domains. Zhang et al. [292] revealed that the blending sequence significantly influences the morphology and electrical properties of the PA6/PP/MWNTs (20/80/4) composites. When blending 3.9 phr MWNTs with a pre-mixed PA6/PP/MWNTs (20/80/0.1) composite, the dispersed PA6 phase forms an elongated structure, which is beneficial to the electrical properties. The same authors [65] also showed that the composite prepared by blending PA6/organoclay (66.7/4) composite with PA6/EPDM-g-MA (13.3/20) blend exhibits the highest impact strength.

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Processing Parameters

The processing parameters generally play an important role in morphology development during polymer blending, especially at higher dispersed phase concentrations and at high values of ηr . The TSE exhibits more efficient dispersive and distributive mixings and higher shear stress than the internal mixer and SSE. Favis and Therrien [276] showed that the phase size of 5% PC in PP is four times larger in an internal mixer than in a TSE at a higher value of ηr (17.3), whereas at a lower ηr the phase size and size distribution are identical in the two mixers. Huang et al. [348, 357], using elastomers (SEBS, EPR and EOR) to toughen the PAs, found that the TSE yields smaller particles with a more narrow distribution of sizes than does the SSE for both the binary and ternary blends as might be expected. However, the choice of extruder does not seem to affect whether there is a bimodal particle size distribution or not. Taylor’s theory (Eqs. (6.1) and (6.2)) showed that the droplet size is inversely proportional to the applied shear stress. This indeed was verified by some researchers. A number of different results, however, have been observed. For PP/PC blend with 7 vol% PC, there is no significant change in dispersed phase diameter over a two- to three-fold increase in the shear stress, regardless of whether an internal mixer [87] or a TSE [205] was used. The morphologies of PC/PP blends with 5% of PP and ηr below 0.22 are not sensitive to the screw speed and output [276]. The dispersed phase size developed during blending may increase, decrease, or show complex non-monotonic behavior when increasing the shear rate [103, 180, 185] due to the competing effects of increased particle–particle contacts versus decreased contact times. Bordereau et al. [358], examining the blending of PS and HDPE in a TSE, found that the domain sizes decrease only slightly (∼25%) over a two-fold increase in the screw speed. In studying the 2/1 LDPE/PS blend, Plochocki et al. [95] indicated that the domain size goes through a minimum as mixing energy is increased. Forteln´y et al. [191] found that the average size of the dispersed phase (PPO+PS) in 50/50 modified PA/(PPO+PS) blend increases first up to a maximum at a speed of 80 rpm, and then decreases. Both results were explained by a consequence of the dynamic equilibrium between breakup and coalescence of droplets in a highly concentrated blend at the condition of steady mixing. Sundararaj and Macosko [103] also found that there is a critical minimum dispersed phase size when increasing the shear rate. This is attributed to the fact that increasing shear rate decreases the matrix viscosity and increases the dispersed phase elasticity, so that the droplet resists the deformation to a greater extent (Eq. (6.17)). The above results imply that polymer blending is governed primarily by dispersive mixing, which involves rupture and distribution of the minor component into a matrix phase. Moreover, it is inappropriate to attempt to correlate the resulting drop size with an average measure of the shear stress in a complex flow field. Polymer nanocomposites have been prepared using a variety of melt processing devices (e.g. extruders, mixers, ultrasonicators, etc.), among which TSEs have proven to be most effective for exfoliation and dispersion [58]. Processing parameters play a key role in achieving a high degree of exfoliation and dispersion of silicate layers in nanocomposites prepared by melt-blending. The better results can be obtained using processing parameters that maximize the shear stress exerted on the polymer matrix. Most studies observed that increasing shear stress or screw speed leads to a better exfoliation and dispersion of clay layer silicates in PP [352, 359–362] and PA6 [81] matrices. This can be attributed to the fact that a higher shear rate allows to break the agglomerates in smaller aggregates. For the effects of other processing parameters, the results reported in the literature are relatively controversial. Lertwimolnun and Vergnes [361] and Nassar et al. [320] found that feed rate has no effect on the state of intercalation. However, decreasing feed rate can significantly improve the state of exfoliation, which is assumed to be related to the corresponding increase in mixing time [55, 361]. Read et al. [363] also suggested that the final dispersion level at the smallest scale for PP/nanoclay nanocomposite is reached in the first melting section of a co-rotating TSE but is not reduced with additional shear or mixing time. On a larger length scale, the number of large agglomerates is reduced with longer mixing time. However, Yang and Ozisik [59] showed that longer mixing time at constant shear stress does not alter the

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clay structure. No clear effect was observed of the mixing temperature on the microstructure of PA6/nanoclay nanocomposite [81], whereas for PP/organoclay nanocomposites, decreasing mixing temperature improves clay layer silicate exfoliation [359, 360]. In fact, it should be noted that as temperature increases, diffusion is improved, which favors to intercalate and exfoliate the platelets. At the same time, the viscosity decreases, and thus the stress necessary to break the clay aggregates decreases. Focusing on the effects of the type of extruder and its screw configurations on the degree of exfoliation and dispersion of layered silicate, Dennis et al. [55] and Cho and Paul [81] extensively prepared PA6/organoclay nanocomposites using intermeshing co-rotating TSE, non-intermeshing counter-rotating TSE, and SSE. It is demonstrated that the degree of exfoliation and dispersion is affected by both the clay chemical treatment and the type and screw configuration of extruders. The non-intermeshing TSE yields the best exfoliation and dispersion. Excellent exfoliation and uniform dispersion can be achieved with both co-rotating and counter-rotating intermeshing extruders when a fully optimized screw configuration is used. Owing to the combination of shear and good polymer–organoclay affinity, TSEs lead to composite properties comparable to those produced by in situ techniques [81], whereas for the composite prepared by the SSE the exfoliation of the clay platelets is not extensive. Even after a second pass through this extruder, undispersed tactoids are still easily observed by the naked eye. This is due to insufficient amount of shear and too short residence time. Increasing the residence time in the extruder generally improves the exfoliation and dispersion. However, there appears to be an optimum extent of backmixing and optimum shear intensity; excessive shear intensity or backmixing apparently causes poorer exfoliation and dispersion [55]. A mechanism for exfoliation and dispersion, as described in Figure 6.12a (Section 6.2.4), should be considered when designing a screw configuration. Shear intensity is required to start the dispersion process, by shearing particles apart into tactoids or intercalants. Residence time in a low or mild shear is required to allow polymer chains to enter the clay galleries and peel the platelets apart. Moreover, the clay could be exfoliated even at low shear rate when the mixing temperatures are high, due to the fact that the diffusion of polymers into the interlayers is enhanced. Interfacial adhesion needs to be higher to improve clay dispersion. However, other researchers have reported on nanocomposite preparations using SSEs. For example, McNally et al. [364] successfully prepared PA12/clay nanocomposites using conventional single-screw meltblending. Often, special screw designs, including provisions for additional mixing, or static mixers at the end of the screw are used to enhance mixing and thus the silicate dispersion. Apart from the various extrusion systems, internal mixers may also be successfully used for the preparation of exfoliated nanocomposites, as demonstrated, for example, in the case of PEI matrix [365]. However, these mixers appear to be much less popular in nanocomposite preparation. Hobbie et al. [366] optically measured the orientation of MWNTs in polymer melts during shear. For a weakly elastic polymer melt, the semiflexible tubes orient along the flow direction at low shear stress, with a transition to vorticity alignment above a critical shear stress. In contrast, for a highly elastic polymer solution, the tubes orient with the flow field at high shear rates.

Nomenclature A: interfacial area AA: acrylic acid ABS: acrylonitrile-butadiene-styrene copolymer AFM: atomic force microscope aPA: amorphous polyamide ASA: acrylonitrile-styrene-acrylic terpolymer

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BEM: boundary element method BIM: boundary integral method C: coalescence efficiency cφ : volume concentration ratio Ca: capillary number CaE : elastic capillary number Cacrit : critical capillary number Cag : global capillary number CB: carbon black CCAB: continuous chaotic advection blender CS: Chaos Screw D: droplet diameter D0 : initial droplet diameter Dn : number-average domain diameter De: Deborah number E-MA-GMA: ethane-(methyl acrylate)-(glycidyl methacrylate) terpolymer EDK : bulk breaking energy EA: ethylene-acrylate copolymer EFM: extensional flow mixer EMA: ethylene methylacrylate ENR: epoxidized natural rubber EO-g-MA: ethane-octene copolymer grafted with maleic anhydride EOR: ethylene/1-octene rubber EOR-g-MA: ethylene/1-octene rubber grafted with maleic anhydride EPDM: ethylene-propylene-diene metallocene terpolymer EPDM-g-MA: ethylene-propylene-diene metallocene terpolymer grafted with maleic anhydride EPM: ethylene-propylene-maleic copolymer EPM-g-MA: ethylene-propylene-maleic copolymer grafted with maleic anhydride EPMA: ethylene-propylene-maleic-anhydride EPR: ethylene-propylene rubber EPR-g-MA: ethylene-propylene rubber grafted with maleic anhydride EVA: ethylene-vinyl acetate copolymer EVA-g-MA: ethylene-vinyl acetate copolymer grafted with maleic anhydride EVOH: ethylene vinyl alcohol copolymer G: Gibbs free energy G  : elastic modulus GMA: glycidyl methacrylate hc : critical separation distance HDPE: high-density polyethylene HMW: high molecular weight I r : interfacial tension ratio IPB: interpenetrating blend iPP: isotactic polypropylene LAOS: large amplitude oscillatory shear flow LDPE: low-density polyethylene LLDPE: linear low-density polyethylene LMW: low molecular weight

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MA: maleic anhydride mEPDM: maleinized ethylene-propylene-diene metallocene terpolymer MMW: medium molecular weight mPE: ethylene-octene copolymer mSEBS: maleinized styrene-ethylene/butylene-styrene triblock copolymer MW: molecular weight MWNTs: multi-walled carbon nanotubes N: numbers of chaotic mixing periods n: number of moles N 1 : first normal stress difference N 2 : second normal stress difference ni : particle numbers of phase i NR: natural rubber Pr : probability that droplets coalesce after their collision PA: polyamide PA11: polyamide 11 PA12: polyamide 12 PA6: polyamide 6 PA66: polyamide 66 PB: polybutylene PBd: polybutadiene PBSA: poly[(butylene succinate)-co-adipate] PBT: poly(butylene terephthalate) PC: polycarbonate PCL: poly(ε-caprolactone) PDMS: poly(dimethyl siloxane) PE: polyethylene PE-g-MA: polyethylene grafted with maleic anhydride Pes : surface P´eclet number PEG: poly(ethylene glycol) PEI: poly(etherimide) PEMA: poly(ethyl methacrylate) PET: poly(ethylene terephthalate) PETG: glycol-modified poly(ethylene terephthalate) PIB: polyisobutylene PLA: polylactide PLLA: poly(L-lactide) PMMA: poly(methyl methacrylate) PO: polyolefine POM: polyoxymethylene PP: polypropylene PP-g-AA: polypropylene grafted with acrylic acid PP-g-GMA: polypropylene grafted with glycidyl methacrylate PP-g-MA: polypropylene grafted with maleic anhydride PP-o-g-MA: polypropylene oligomer grafted with maleic anhydride PPO: poly(phenylene oxide) PS: polystyrene

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PS-g-MA: polystyrene grafted with maleic anhydride PSoz: poly(styrene-co-methyl vinyl oxazoline) PVC: poly(vinyl chloride) PVDF: poly(vinylidene fluoride) Sr : stress ratio SAN: styrene-acrylonitrile copolymer SAN-g-MA: styrene-acrylonitrile copolymer grafted with maleic anhydride SBR: styrene butadiene rubber SEBS: styrene-ethylene/butylene-styrene triblock copolymer SEBS-g-MA: styrene-ethylene/butylene-styrene triblock copolymer grafted with maleic anhydride SEM: scanning electron microscope SSE: single-screw extruder SSSP: solid-state shear pulverization T 11 : first normal stress T 22 : second normal stress tc : critical time tcc : critical coalescence time TEM: transmission electron microscope TPO: thermoplastic olefin TPU: thermoplastic polyurethane TSE: twin-screw extruder V i : volume fraction of phase i Wi: Weissenberg number WAXD: wide-angle X-ray diffraction γ : shear strain γ˙ : shear rate ε˙ : strain rate σ : interfacial stress φ: volume fraction η: viscosity ηr : viscosity ratio η+ E,m : transient elongational viscosity of matrix λ: relaxation time λ L : Lyapunov exponent λij : spreading coefficient of i over j μ: chemical potential Subscripts d: dispersed phase or droplet m: matrix phase

Acknowledgements Financial support provided by the National Natural Science Foundation of China (20874031, 10672061) is gratefully acknowledged.

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References 1. R.W. Flumerfelt, Drop breakup in simple shear fields of viscoelastic fluids, Ind. Eng. Chem. Fundam., 11, 312–318 (1972). 2. H. VanOene, Modes of dispersion of viscoelastic fluids in flow, J. Colloid Interface Sci., 40, 448–467 (1972). 3. S. Wu, Formation of dispersed phase in incompatible polymer blends: interfacial and rheological effects, Polym. Eng. Sci., 27, 335–343 (1987). 4. L. Levitt, C.W. Macosko, S.D. Pearson, Influence of normal stress difference on polymer drop deformation, Polym. Eng. Sci., 36, 1647–1655 (1996). 5. F. Mighri, A. Ajji, P.J. Carreau, Influence of elastic properties on drop deformation in elongational flow, J. Rheol., 41, 1183–1201 (1997). 6. F. Mighri, P.J. Carreau, A. Ajji, Influence of elastic properties on drop deformation and breakup in shear flow, J. Rheol., 42, 1477–1490 (1998). 7. I. Luzinov, C. Pagnoulle, R. J´erˆome, Dependence of phase morphology and mechanical properties of PS/SBR/PE ternary blends on composition: transition from core–shell to triple-phase continuity structures, Polymer, 41, 3381–3389 (2000). 8. F. Mighri, M.A. Huneault, Dispersion visualization of model fluids in a transparent Couette flow cell, J. Rheol., 45, 783–797 (2001). 9. W. Lerdwijitjarud, A. Sirivat, R.G. Larson, Influence of elasticity on dispersed-phase droplet size in immiscible polymer blends in simple shearing flow, Polym. Eng. Sci., 42, 798–809 (2002). 10. P. Moldenaers, C.L. Tucker, Microstructural evolution in polymer blends, Annu. Rev. Fluid Mech., 34, 177–210 (2002). 11. B. Lin, F. Mighri, M.A. Huneault, U. Sundararaj, Parallel breakup of polymer drops under simple shear, Macromol. Rapid Commun., 24, 783–788 (2003). 12. B. Lin, U. Sundararaj, F. Mighri, M.A. Huneault, Erosion and breakup of polymer drops under simple shear in high viscosity ratio systems, Polym. Eng. Sci., 43, 891–904 (2003). 13. W. Lerdwijitjarud, A. Sirivat, R.G. Larson, Influence of dispersed-phase elasticity on steady-state deformation and breakup of droplets in simple shearing flow of immiscible polymer blends, J. Rheol., 48, 843–862 (2004). 14. B. Li, U. Sundararaj, Sheet formation during drop deformation and breakup in polyethylene/polycarbonate systems sheared between parallel plates, Polymer, 45, 7605–7613 (2004). 15. B. Lin, F. Mighri, M.A. Huneault, U. Sundararaj, Effect of premade compatibilizer and reactive polymers on polystyrene drop deformation and breakup in simple shear, Macromolecules, 38, 5609–5616 (2005). 16. H. Chen, U. Sundararaj, K. Nandakumar, M.D. Wetzel, Erosion of polymer pellets during blending in a twin-screw extruder, AIChE J., 52, 1267–1270 (2006). 17. F. Mighri, M.A. Huneault, In situ visualization of drop deformation, erosion, and breakup in high viscosity ratio polymeric systems under high shearing stress conditions, J. Appl. Polym. Sci., 100, 2582–2591 (2006). 18. H.P. Li, U. Sundararaj, Morphology development of polymer blends in extruder: the effects of compatibilization and rotation rate, Macromol. Chem. Phys., 210, 852–863 (2009). 19. P. P¨otschke, D.R. Paul, Formation of co-continuous structures in melt-mixed immiscible polymer blends, J. Macromol. Sci.-Polym. R., C43, 87–141 (2003). 20. H.X. Huang, G. Jiang, X.J. Li, Development of polymer blend morphology along an extruder with different screw geometries, Inter. Polym. Process., 23, 47–54 (2008). 21. H.X. Li, G.H. Hu, The early stage of the morphology development of immiscible polymer blends during melt blending: compatibilized vs. uncompatibilized blends, J. Polym. Sci. Part B: Polym. Phys., 39, 601–610 (2001). 22. L.A. Utracki, Compatibilization of polymer blends, Can. J. Chem. Eng., 80, 1008–1015 (2002). 23. C.W. Macosko, Morphology development and control in immiscible polymer blends, Macromol. Symp., 149, 171–184 (2000). 24. R.C. Willemse, A. Posthuma de Boer, J. van Dam, A.D. Gotsis, Co-continuous morphologies in polymer blends: a new model, Polymer, 39, 5879–5887 (1998). 25. S. Wu, Phase structure and adhesion in polymer blends: a criterion for rubber toughening, Polymer, 26, 1855–1863 (1985).

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26. C.C. Chen, J.L. White, Compatibilizing agents in polymer blends: interfacial tension, phase morphology, and mechanical properties, Polym. Eng. Sci., 33, 923–930 (1993). 27. K.B. Migler, E.K. Hobbie, F. Qiao, In line study of droplet deformation in polymer blends in channel flow, Polym. Eng. Sci., 39, 2282–2291 (1999). 28. L. Barangi, H. Nazockdast, F. Afshar Taromi, Effect of the melt viscoelastic behavior of components on the morphology development of polymer blends in a twin-screw extruder, J. Appl. Polym. Sci., 108, 2558–2563 (2008). 29. I. Vinckier, P. Moldenaers, J. Mewis, Relationship between rheology and morphology of model blends in steady shear flow, J. Rheol., 40, 613–631 (1996). 30. C.E. Scott, C.W. Macosko, Morphology development during the initial stages of polymer-polymer blending, Polymer, 36, 461–470 (1995). 31. K. Verhulst, R. Cardinaels, P. Moldenaers, S. Afkhami, Y. Renardy, Influence of viscoelasticity on drop deformation and orientation in shear flow. Part 2: dynamics, J. Non-Newtonian Fluid Mech., 156, 44–57 (2009). 32. B. Lin, U. Sundararaj, P. Guegan, Effect of mixing protocol on compatibilized polymer blend morphology, Polym. Eng. Sci., 46, 691–702 (2006). 33. H.F. Guo, S. Packirisamy, N.V. Gvozdic, D.J. Meier, Prediction and manipulation of the phase morphologies of multiphase polymer blends: 1. ternary systems, Polymer, 38, 785–794 (1997). 34. I. Luzinov, K. Xi, C. Pagnoulle, G. Huynh-Ba, R. J´erˆome, Composition effect on the core–shell morphology and mechanical properties of ternary polystyrene/styrene–butadiene rubber/polyethylene blends, Polymer, 40, 2511–2520 (1999). 35. A.K. Gupta, K.R. Srinivasan, Melt rheology and morphology of PP/SEBS/PC ternary blend, J. Appl. Polym. Sci., 47, 167–184 (1993). 36. K.P. Tchomakov, B.D. Favis, M.A. Huneault, M.F. Champagne, F. Tofan, Mechanical properties and morphology of ternary PP/EPDM/PE blends, Can. J. Chem. Eng., 83, 300–309 (2005). 37. S.Y. Hobbs, M.E.J. Dekkers, V.H. Watkins, Effect of interfacial forces on polymer blend morphologies, Polymer, 29, 1598–1602 (1988). 38. K.P. Tchomakov, B.D. Favis, M.A. Huneault, M.F. Champagne, F. Tofan, Composite droplets with core/shell morphologies prepared from HDPE/PS/PMMA ternary blends by twin-screw extrusion, Polym. Eng. Sci., 44, 749–759 (2004). 39. C. Harrats, S. Thomas, G. Groeninckx, Micro- and nanostructured multiphase polymer blend systems: phase morphology and interfaces, Taylor & Francis, New York, 2006. 40. G.H. Hu, H. Cartier, Reactive extrusion: toward nanoblends, Macromolecules, 32, 4713–4718 (1999). 41. H. Pernot, M. Baumert, F. Court, L. Leibler, Design and properties of co-continuous nanostructured polymers by reactive blending, Nat. Mater., 1, 54–58 (2002). 42. A.J. Busby, J.X. Zhang, A. Naylor, C.J. Roberts, M.C. Davies, S.J.B. Tendler, et al., The preparation of novel nanostructured polymer blends of ultra high molecular weight polyethylene with polymethacrylates using supercritical carbon dioxide, J. Mater. Chem., 13, 2838–2844 (2003). 43. G.H. Hu, L.F. Feng, Extruder processing for nanoblends and nanocomposites, Macromol. Symp., 195, 303–308 (2003). 44. L. Gani, S. Tenc´e-Girault, M. Mill´equant, S. Bizet, L. Leibler, Co-continuous nanostructured blend by reactive blending: incorporation of high molecular weight polymers, Macromol. Chem. Phys., 211, 736–743 (2010). 45. H. Shimizu, Y.J. Li, A. Kaito, H. Sano, Formation of nanostructured PVDF/PA11 blends using high-shear processing, Macromolecules, 38, 7880–7883 (2005). 46. Y. Tao, J. Kim, J.M. Torkelson, Achievement of quasi-nanostructured polymer blends by solid-state shear pulverization and compatibilization by gradient copolymer addition, Polymer, 47, 6773–6781 (2006). 47. P.M. Lipic, F.S. Bates, M.A. Hillmyer, Nanostructured thermosets from self-assembled amphiphilic block copolymer/epoxy resin mixtures, J. Am. Chem. Soc., 120, 8963–8970 (1998). 48. M.A. Amendt, L. Chen., M.A. Hillmyer, Formation of nanostructured poly(dicyclopentadiene) thermosets using reactive block polymers, Macromolecules, 43, 3924–3934 (2010). 49. M. Alexandre, P. Dubois, Polymer-layered silicate nanocomposites: preparation, properties and uses of a new class of materials. Mater. Sci. Eng. R., 28, 1–63 (2000).

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50. S. Pavlidou, C.D. Papaspyrides, A review on polymer–layered silicate nanocomposites, Prog. Polym. Sci., 33, 1119–1198 (2008). 51. T. Kietzke, D. Neher, K. Landfester, R. Montenegro, R. Guntner, U. Scherf, Novel approaches to polymer blends based on polymer nanoparticles, Nat. Mater., 2, 408–412 (2003). 52. L.W. Carter, J.G. Hendricks, D.S. Bolley, Elastomer reinforced with modified clay, US Patent, 2,531,396 (1950). 53. A. Dasari, Z.-Z. Yu, Y.-W. Mai, Effect of blending sequence on microstructure of ternary nanocomposites, Polymer, 46: 5986–5991 (2005). 54. S.S. Ray, M. Okamoto, Polymer/layered silicate nanocomposite: a review from preparation to processing, Prog. Polym. Sci., 28, 1539–1641 (2003). 55. H.R. Dennis, D.L. Hunter, D. Chang, S. Kim, J.L. White, J.W. Cho, D.R. Paul, Effect of melt processing conditions on the extent of exfoliation in organoclay-based nanocomposites, Polymer, 42, 9513–9522 (2001). 56. F. Hussain, M. Hojjati, M. Okamoto, R.E. Gorga, Polymer-matrix nanocomposites, processing, manufacturing, and application: an overview, J. Compos. Maters., 40, 1511–1575 (2006). 57. M.A. Treece, W. Zhang, R.D. Moffitt, J.P. Oberhauser, Twin-screw extrusion of polypropylene-clay nanocomposites: influence of masterbatch processing, screw rotation mode, and sequence, Polym. Eng. Sci., 47, 898–911 (2007). 58. T.D. Fornes, P.J. Yoon, H. Keskkula, D.R. Paul, Nylon 6 nanocomposites: the effect of matrix molecular weight, Polymer, 42, 9929–9940 (2001). 59. K. Yang, R. Ozisik, Effects of processing parameters on the preparation of nylon 6 nanocomposites, Polymer, 47, 2849–2855 (2006). 60. T.D. Fornes, P.J. Yoon, D.L. Hunter, H. Keskkula, D.R. Paul, Effect of organoclay structure on nylon-6 nanocomposite morphology and properties, Polymer, 43, 5915–5933 (2002). 61. B.B. Khatua, D.J. Lee, H.Y. Kim, J.K. Kim, Effect of organoclay platelets on morphology of nylon-6 and poly(ethylene-ran-propylene) rubber blends, Macromolecules, 37, 2454–2459 (2004). 62. Y. Yoo, C. Park, S.G. Lee, K.Y. Choi, D.S. Kim, J.H. Lee, Influence of addition of organoclays on morphologies in nylon 6/LLDPE blends, Macromol. Chem. Phys., 206, 878–884 (2005). 63. F.C. Chiu, S.M. Lai, Y.L. Chen, T.H. Lee, Investigation on the polyamide 6/organoclay nanocomposites with or without a maleated polyolefin elastomer as a toughener, Polymer, 46, 11600–11609 (2005). 64. Y.C. Ahn, D.R. Paul, Rubber toughening of nylon 6 nanocomposites, Polymer, 47, 2830–2838 (2006). 65. L.Y. Zhang, C.Y. Wan, Y. Zhang, Investigation on morphology and mechanical properties of polyamide 6/maleated ethylene-propylene-diene rubber/organoclay composites, Polym. Eng. Sci., 49, 209–216 (2009). 66. S.S. Ray, M. Bousmina, Effect of organic modification on the compatibilization efficiency of clay in an immiscible polymer blend, Macromol. Rapid Commun., 26, 1639–1646 (2005). 67. S.S. Ray, S. Pouliot, M. Bousmina, L.A. Utracki, Role of organically modified layered silicate as an active interfacial modifier in immiscible polystyrene/polypropylene blends, Polymer, 45, 8403–8413 (2004). 68. S.S. Ray, J. Bandyopadhyay, M. Bousmina, Effect of organoclay on the morphology and properties of poly(propylene)/poly[(butylene succinate)-co-adipate] blends, Macromol. Mater. Eng., 292, 729–747 (2007). 69. M.Y. Gelfer, H.H. Song, L.Z. Liu, B.S. Hsiao, B. Chu, M. Rafailovich, M.Y. Si, V. Zaitsev, Effects of organoclays on morphology and thermal and rheological properties of polystyrene and poly(methyl methacrylate) blends, J. Polym. Sci. Part B: Polym. Phys., 41, 44–54 (2003). 70. Y.J. Li, H. Shimizu, Novel morphologies of poly(phenylene oxide) (PPO)/polyamide 6 (PA6) blend nanocomposites, Polymer, 45, 7381–7388 (2004). 71. Y.J. Li, H. Shimizu, Co-continuous polyamide 6 (PA6)/acrylonitrile-butadiene-styrene (ABS) nanocomposites, Macromol. Rapid Commun., 26, 710–715 (2005). 72. M. Gahleitner, B. Kretzschmar, V.G. Vliet, J. Devaux, D. Pospiech, K. Bernreitner, E. Ingolic, Rheology/morphology interactions in polypropylene/polyamide-6 nanocomposites, Rheol. Acta., 45, 322–330 (2006). 73. D. Voulgaris, D. Petridis, Emulsifying effect of dimethyldioctadecylammonium-hectorite in polystyrene/poly(ethyl methacrylate) blends, Polymer, 43, 2213–2218 (2002). 74. D.F. Wu, C.X. Zhou, M. Zhang, Effect of clay on immiscible morphology of poly(butylene terephthalate)/polyethylene blend nanocomposites, J. Appl. Polym. Sci., 102, 3628–3633 (2006). 75. J. S. Hong, Y. K. Kim, K. H. Ahn, S. J. Lee, C. Kim, Interfacial tension reduction in PBT/PE/clay nanocomposite, Rheol. Acta., 46, 469–478 (2007).

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76. R.A. Vaia, H. Ishii, E.P. Giannelis, Synthesis and properties of two-dimensional nanostructures by direct intercalation of polymer melts in layered silicates, Chem. Mater., 5, 1694–1696 (1993). 77. R.A. Vaia, E.P. Giannelis, Lattice of polymer melt intercalation in organically modified layered silicates, Macromolecules, 30, 7990–7999 (1997). 78. R.A. Vaia, E.P. Giannelis, Polymer melt intercalation in organically-modified layered silicates: model predictions and experiment, Macromolecules, 30, 8000–8009 (1997). 79. M. Kawasumi, N. Hasegawa, M. Kato, A. Usuki, A. Okada, Preparation and mechanical properties of polypropyleneclay hybrids, Macromolecules, 30, 6333–6338 (1997). 80. N. Hasegawa, H. Okamoto, M. Kawasumi, M. Kato, A. Tsukigase, A. Usuki, Polyolefin-clay hybrids based on modified polyolefins and organoclay, Macromol. Mater. Eng., 280/281, 76–79 (2000). 81. J.W. Cho, D.R. Paul, Nylon 6 nanocomposites by melt compounding, Polymer, 42, 1083–1094 (2001). 82. G. Filippone, N.Tz. Dintcheva, D. Acierno, F.P. La Mantia, The role of organoclay in promoting co-continuous morphology in high-density poly(ethylene)/poly(amide) 6 blends, Polymer, 49, 1312–1322 (2008). 83. M. Kontopoulou, Y.Q. Liu, J.R. Austin, J.S. Parent, The dynamics of montmorillonite clay dispersion and morphology development in immiscible ethylene-propylene rubber/polypropylene blends, Polymer, 48, 4520–4528 (2007). 84. K. Li, H.X. Huang, G. Jiang, Comparing the continuous chaotic and shear mixing-induced morphology development of polyblend and its nanocomposites, Polym. Plast. Technol. Eng., 48, 989–995 (2009). 85. C.E. Scott, C.W. Macosko, Model experiments concerning morphology development during the initial stages of polymer blending, Polym. Bull., 26, 341–348 (1991). 86. U. Sundararaj, C.W. Macosko, C.K. Shih, Evidence for inversion of phase continuity during morphology development in polymer blending, Polym. Eng. Sci., 36, 1769–1781 (1996). 87. B.D. Favis, The effect of processing parameters on the morphology of an immiscible binary blend, J. Appl. Polym. Sci., 39, 285–300 (1990). 88. X. Chen, J. Xu, B.H. Guo, Development of dispersed phase size and its dependence on processing parameters, J. Appl. Polym. Sci., 102, 3201–3211 (2006). 89. U. Sundararaj, C.W. Macosko, R.J. Rolando, H.T. Chan, Morphology development in polymer blends, Polym. Eng. Sci., 32, 1814–1823 (1992). 90. H. Potente, M. Bastian, K. Bergemann, M. Senge, G. Scheel, Th. Winkelmann, Morphology of polymer blends in the melting section of co-rotating twin screw extruders, Polym. Eng. Sci., 41, 222–231 (2001). 91. J.T. Lindt, A.K. Ghosh, Fluid mechanics of the formation of polymer blends. Part I: formation of lamellar structures, Polym. Eng. Sci., 32, 1802–1813 (1992). 92. H.E. Burch, C.E. Scott, Effect of viscosity ratio on structure evolution in miscible polymer blends, Polymer, 42, 7313–7325 (2001). 93. J. Leblanc, F. Pla, M. Mercier, P. Pitiot, Elaboration of PEG/PA66 blends in a twin-screw batch-type mini-extruder: process study and modeling, Chem. Eng. Sci., 64, 1918–1926 (2009). 94. J. Reignier, B.D. Favis, Control of the subinclusion microstructure in HDPE/PS/PMMA ternary blends, Macromolecules, 33, 6998–7008 (2000). 95. A.P. Plochocki, S.S. Dagli, R.D. Andrews, The interface in binary mixtures of polymers containing a corresponding block copolymer: effects of industrial mixing processes and of coalescence, Polym. Eng. Sci., 30, 741–752 (1990). 96. F. Mighri, M.A. Huneault, Drop deformation and breakup mechanisms in viscoelastic model fluid systems and polymer blends, Can J. Chem. Eng., 80, 1028–1035 (2002). 97. S.P. Rwei, I. Manas-Zloczower, D.L. Feke, Characterization of agglomerate dispersion by erosion in simple shear flows, Polym. Eng. Sci., 31, 558–562 (1991). 98. H. Chen, U. Sundararaj, K. Nandakumar, Modeling of polymer melting, drop deformation and breakup under shear flow, Polym. Eng. Sci., 44, 1258–1266 (2004). 99. K. Min, J.L. White, Rubber Chem. Technol., 58, 1024 (1985). 100. U. Sundararaj, C.W. Macosko, A. Nakayama, T. Inoue, Milligrams to kilograms: an evaluation of mixers for reactive polymer blending, Polym. Eng. Sci., 35, 100–114 (1995). 101. U. Sundararaj, Y. Dori, C.W. Macosko, Sheet formation in immiscible polymer blends: model experiments on initial blend morphology, Polymer, 36, 1957–1968 (1995).

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102. R.C. Willemse, E.J.J. Ramaker, J. van Dam, A. Posthuma de Boer, Morphology development in immiscible polymer blends: initial blend morphology and phase dimensions, Polymer, 40, 6651–6659 (1999). 103. U. Sundararaj, C.W. Macosko, Drop breakup and coalescence in polymer blends: the effects of concentration and compatibilization, Macromolecules, 28, 2647–2657 (1995). 104. P.G. Ghodgaonkar, U. Sundararaj, Prediction of dispersed phase drop diameter in polymer blends: the effect of elasticity, Polym. Eng. Sci., 36, 1656–1665 (1996). 105. L.A. Utracki, Z.H. Shi, Development of polymer blend morphology during compounding in a twin-screw extruder. Part I: Droplet dispersion and coalescence—a review, Polym. Eng. Sci., 32, 1824–1833 (1992). 106. W. Lerdwijitjarud, R.G. Larson, A. Sirivat, M.J. Solomon, Influence of weak elasticity of dispersed phase on droplet behavior in sheared polybutadiene/poly(dimethyl siloxane) blends, J. Rheol., 47, 37–58 (2003). 107. T. Cherdhirankorn, W. Lerdwijitjarud, A. Sirivat, R. G. Larson, Dynamics of vorticity stretching and breakup of isolated viscoelastic droplets in an immiscible viscoelastic matrix, Rheol. Acta., 43, 246–256 (2004). 108. K.B. Migler, Droplet vorticity alignment in model polymer blends, J. Rheol., 44, 277–290 (2000). 109. K.B. Migler, String formation in sheared polymer blends: coalescence, breakup, and finite size effects, Phys. Rev. Lett., 86, 1023–1026 (2001). 110. P.P. Varanasi, M.E. Ryan, P. Stroeve, Experimental study on the breakup of model viscoelastic drops in uniform shear flow, Ind. Eng. Chem., 33, 1858–1866 (1994). 111. V.T. Tsakalos, P. Navard, E. Peuvrel-Disdier, Deformation and breakup mechanisms of single drops during shear, J. Rheol., 42, 1403–1417 (1998). 112. S. Guido, M. Villone, Three-dimensional shape of a drop under simple shear flow, J. Rheol., 42, 395–415 (1998). 113. S. Guido, M. Simeone, F. Greco, Deformation of a Newtonian drop in a viscoelastic matrix under steady shear flow: Experimental validation of slow flow theory, J. Non-Newtonian Fluid Mech., 114, 65–82 (2003). 114. V. Sibillo, M. Simeone, S. Guido, Break-up of a Newtonian drop in a viscoelastic matrix under simple shear flow, Rheol. Acta., 43, 449–456 (2004). 115. G.I. Taylor, The deformation of emulsions in definable fields of flow, Proc. R. Soc. London, A146, 501–523 (1934). 116. B.H. Birkhofer, J.C. Eischen, D. Megias-Alguacil, P. Fischer, E.J. Windhab, Computer-controlled flow cell for the study of particle and drop dynamics in shear flow fields, Ind. Eng. Chem. Res., 44, 6999–7009 (2005). 117. H.P. Li, U. Sundararaj, Does drop size affect the mechanism of viscoelastic drop breakup?, Phys. Fluids, 20, 053101 (2008). 118. D. Bourry, F. Godbille, R.E. Khayat, A. Luciani, J. Picot, L.A. Utracki, Extensional flow of polymeric dispersions, Polym. Eng. Sci., 39, 1072–1086 (1999). 119. R.E. Khayat, A. Luciani, L.A. Utracki, F. Godbille, J. Picot, Influence of shear and elongation on drop deformation in convergent-divergent flows, Inter. J. Multiphase Flow, 26, 17–44 (2000). 120. F.D. Godbille, J.J.C. Picot, Drop breakup in combined shear and extensional flow conditions, Adv. Polym. Technol., 19, 14–21 (2000). 121. M. Meller, A. Luciani, J.-A. E. M˚anson, Flow through a convergence. Part 2: mixing of high viscosity ratio polymer blends, Polym. Eng. Sci., 42, 634–653 (2002). 122. Y.T. Hu, D.J. Pine, L.G. Leal, Drop deformation, breakup and coalescence with compatibilizer, Phys. Fluids, 12, 484–489 (2000). 123. J.W. Ha, L.G. Leal, An experimental study of drop deformation and breakup in extensional flow at high capillary number, Phys. Fluids, 13, 1568–1576 (2001). 124. D.C. Tretheway, L.G. Leal, Deformation and relaxation of Newtonian drops in planar extensional flows of a Boger fluid, J. Non-Newtonian Fluid Mech., 99, 81–108 (2001). 125. U.A. Handge, Experimental investigation of equibiaxial extension and breakup of drops in a molten two-phase polymer blend, Phys. Rev. E., 72, 011801 (2005). 126. A.S. Hsu, L.G. Leal, Deformation of a viscoelastic drop in planar extensional flows of a Newtonian fluid, J. NonNewtonian Fluid Mech., 160, 176–180 (2009). 127. J.M. Ottino, C.M. Leong, H. Rising, P.D. Swanson, Morphological structures produced by mixing in chaotic flows, Nature, 333, 419–425 (1988). 128. M. Tjahjadi, J.M. Ottino, Stretching and breakup of droplets in chaotic flows, J. Fluid Mech., 232, 191–219 (1991).

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129. H.A. Kusch, J.M. Ottino, Experiments on mixing in continuous chaotic flows, J. Fluid Mech., 236, 319–348 (1992). 130. C.A. Florek, C.L. Tucker, Stretching distributions in chaotic mixing of droplet dispersions with unequal viscosities, Phys. Fluids, 17, 053101 (2005). 131. T.N. Pham, C.L. Tucker, Stretch and shape distributions of droplets with interfacial tension in chaotic mixing, Inter. Polym. Process., 20, 128–135 (2005). 132. E. Boonen, P. Van Puyvelde, P. Moldenaers, Droplet dynamics in sub-critical complex flows, Rheol. Acta., 48, 359–371 (2009). 133. G.I. Taylor, The viscosity of a fluid containing small drops of another fluid, Proc. R. Soc. London, Ser. A 138: 41–48 (1932). 134. H.P. Grace, Dispersion phenomena in high viscosity immiscible fluid systems and application of static mixers as dispersion devices in such systems, Chem. Eng. Commun., 14, 225–227 (1982). 135. J.J. Elmendorp, R.J. Maalcke, A study on polymer blending microrheology: Part 1, Polym. Eng. Sci., 25, 1041–1047 (1985). 136. H.J. Karam, J.C. Bellinger, Deformation and breakup of liquid droplets in a simple shear field, Ind. Eng. Chem. Fundam., 7, 576–581 (1968). 137. D.I. Bigio, C.R. Marks, R.V. Calabrese, Predicting drop breakup in complex flows from model flow experiments, Inter. Polym. Process., 13, 192–198 (1998). 138. V. Cristini, S. Guido, A. Alfani, J. Blawzdziewicz, M. Loewenberg, Drop breakup and fragment size distribution in shear flow, J. Rheol., 47, 1283–1298 (2003). 139. H. Yamane, M. Takahashi, R. Hayashi, K. Okamoto, H. Kashihara, T. Masuda, Observation of deformation and recovery of poly(isobutylene) droplet in a poly(isobutylene)/poly(dimethyl siloxane) blend after application of step shear strain, J. Rheol., 42, 567–580 (1998). 140. R. Hayashi, M. Takahashi, H. Yamane, H. Jinnai, H. Watanabe, Dynamic interfacial properties of polymer blends under large step strains: shape recovery of a single droplet, Polymer, 42, 757–764 (2001). 141. S. Guido, M. Minale, P.L. Maffettone, Drop shape dynamics under shear-flow reversal, J. Rheol., 44, 1385–1399 (2000). 142. R. Hayashi, M. Takahashi, T. Kajihara, H. Yamane, Application of large double-step shear strains to analyze deformation and shape recovery of a polymer droplet in an immiscible polymer matrix, J. Rheol., 45, 627–639 (2001). 143. S. Wannaborworn, M.R. Mackley, Y. Renardy, Experimental observation and matching numerical simulation for the deformation and breakup of immiscible drops in oscillatory shear, J. Rheol., 46, 1279–1293 (2002). 144. R. Cavallo, S. Guido, M. Simeone, Drop deformation under small-amplitude oscillatory shear flow, Rheol. Acta., 42, 1–9 (2003). 145. S. Guido, M. Grosso, P.L. Maffettone, Newtonian drop in a Newtonian matrix subjected to large amplitude oscillatory shear flows, Rheol. Acta., 43, 575–583 (2004). 146. Y. Renardy, Numerical simulation of a drop undergoing large amplitude oscillatory shear, Rheol. Acta., 45, 223–227 (2006). 147. X.Y. Li, K. Sarkar, Drop dynamics in an oscillating extensional flow at finite Reynolds numbers, Phys. Fluids, 17, 027103 (2005). 148. V. Janpaen, S. Niamlang, W. Lerdwijitjarud, A. Sirivat, Oscillatory shear induced droplet deformation and breakup in immiscible polymer blends, Phys. Fluids, 21, 063102 (2009). 149. Y. Deyrail, M.A. Huneault, M. Bousmina, Rheo-optical study of nonlinear effects in an immiscible newtonian polymer blend under large amplitude oscillatory shear flow, J. Polym. Sci. Part B: Polym. Phys., 47, 1467–1480 (2009). 150. K. Feigl, S.F.M. Kaufmann, P. Fischer, E.J. Windhab, A numerical procedure for calculating droplet deformation in dispersing flows and experimental verification, Chem. Eng. Sci., 58, 2351–2363 (2003). 151. Y.W. Kruijt-Stegeman, F.N. van de Vosse, H.E.H. Meijer, Droplet behavior in the presence of insoluble surfactants, Phys. Fluids, 16, 2785–2796 (2004). 152. R.E. Khayat, M. Huneault, L.A. Utracki, R. Duquette, A boundary element analysis of planar drop deformation in the screw channel of a mixing extruder, Eng. Anal. Bound. Elem., 21, 155–168 (1998).

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153. R.E. Khayat, A boundary-only approach to the deformation of a shear-thinning drop in extensional Newtonian flow, Inter. J. Numer. Meth. Fluids, 33, 559–581 (2000). 154. R.E. Khayat, Three-dimensional boundary element analysis of drop deformation in confined flow for Newtonian and viscoelastic systems, Inter. J. Numer. Meth. Fluids, 34, 241–275 (2000). 155. Y.C. Wang, P. Dimitrakopoulos, A three-dimensional spectral boundary element algorithm for interfacial dynamics in Stokes flow, Phys. Fluids, 18, 082106 (2006). 156. K. Feigl, D. Megias-Alguacil, P. Fischer, E.J. Windhab, Simulation and experiments of droplet deformation and orientation in simple shear flow with surfactants, Chem. Eng. Sci., 62, 3242–3258 (2007). 157. Y. Renardy, M. Renardy, PROST: a parabolic reconstruction of surface tension for the volume-of-fluid method, J. Comput. Phys., 183, 400–421 (2002). 158. T. Chinyoka, Y. Renardy, M. Renardy, D.B. Khismatullin, Two-dimensional study of drop deformation under simple shear for Oldroyd-B liquids, J. Non-Newtonian Fluid Mech., 130, 45–56 (2005). 159. D. Khismatullin, Y. Renardy, M. Renardy, Development and implementation of VOF-PROST for 3D viscoelastic liquid–liquid simulations, J. Non-Newtonian Fluid Mech., 140, 120–131 (2006). 160. S. Afkhami, P. Yue, Y. Renardy, A comparison of viscoelastic stress wakes for two-dimensional and threedimensional Newtonian drop deformations in a viscoelastic matrix under shear, Phys. Fluids, 21, 072106 (2009). 161. N. Aggarwal, K. Sarkar, Deformation and breakup of a viscoelastic drop in a Newtonian matrix under steady shear, J. Fluid Mech., 584, 1–21 (2007). 162. M. Muradoglu, G. Tryggvason, A front-tracking method for computation of interfacial flows with soluble surfactants, J. Comp. Phys., 227, 2238–2262 (2008). 163. P. Dimitrakopoulos, J.T. Wang, A spectral boundary element algorithm for interfacial dynamics in two-dimensional Stokes flow based on Hermitian interfacial smoothing, Eng. Anal. Bound. Elem., 31, 646–656 (2007). 164. V. Cristini, J. Blawzdziewicz, M. Loewenberg, Droplet breakup in three-dimensional viscous flows, Phys. Fluids, 10, 1781–1783 (1998). 165. V. Cristini, J. Blawzdziewicz, M. Loewenberg, An adaptive mesh algorithm for evolving surfaces: Simulations of drop breakup and coalescence, J. Computational Phys., 168, 445–463 (2001). 166. P.J.A. Janssen, P.D. Anderson, Boundary-integral method for drop deformation between parallel plates, Phys. Fluids, 19, 043602 (2007). 167. A. Vananroye, P.J.A. Janssen, P.D. Anderson, P. Van Puyvelde, P. Moldenaers, Microconfined equiviscous droplet deformation: comparison of experimental and numerical results, Phys. Fluids, 20, 013101 (2008). 168. P.J.A. Janssen, P.D. Anderson, A boundary-integral model for drop deformation between two parallel plates with non-unit viscosity ratio drops, J. Computational Phys., 227, 8807–8819 (2008). 169. B.D. Favis, Factors influencing the nlorphology of immiscible polvmer blends in melt processing, In Polymer Blends: Formulation and Performance, D.R. Paul, C.B. Bucknall (Eds), John Wiley & Sons, Inc, New York, 2000. 170. S. Tomotika, On the instability of a cylindrical thread of a viscous liquid surrounded by another viscous fluid, Proc. R. Soc. London, A150, 322–337 (1935). 171. J. Eggers, Nonlinear dynamics and breakup of free surface flows, Rev. Mod. Phys., 69, 865–929 (1997). 172. J.M. Rallison, A numerical study of the deformation and burst of a viscous drop in general shear flows, J. Fluid Mech., 109, 465–482 (1981). 173. H.A. Stone, B.J. Bentley, G.L. Leal, An experimental study of transient effects in the breakup of viscous drops, J. Fluid Mech., 173, 131–151 (1986). 174. X.Y. Zhao, Drop breakup in dilute Newtonian emulsions in simple shear flow: new drop breakup mechanisms, J Rheol., 51, 367–392 (2007). 175. W.J. Milliken, L.G. Leal, Deformation and breakup of viscoelastic drops in planar extensional flows, J. NonNewtonian Fluid Mech., 40, 355–379 (1991). 176. W.J. Milliken, L.G. Leal, A note on the effect of vorticity on the deformation and breakup of polymer drops, J. Non-Newtonian Fluid Mech., 42, 231–239 (1992). 177. W.J. Milliken, L.G. Leal, The influence of surfactant on the deformation and breakup of a viscous drop—the effect of surfactant solubility, J. Colloid Sci., 166, 275–285 (1994). 178. D.C. Tretheway, L.G. Leal, Surfactant and viscoelastic effects on drop deformation in 2-D extensional flow, AIChE J., 45, 929–937 (1999).

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179. E.K. Hobbie, K.B. Migler, Vorticity elongation in polymeric emulsions, Phys. Rev. Lett., 82, 5393–5396 (1999). ˇ y, Coalescence in molten quiescent polymer blends, Polymer, 36, 4113–4118 (1995). 180. I. Forteln´y, A. Zivn´ 181. V. Everaert, L. Aerts, G. Groeninckx, Phase morphology development in immiscible PP/(PS/PPE) blends: influence of the melt-viscosity ratio and blend composition, Polymer, 40, 6627–6644 (1999). 182. V.E. Ziegler, B.A. Wolf, Bimodal drop size distributions during the early stages of shear induced coalescence, Polymer, 46, 9265–9273 (2005). 183. N. Tokita, Analysis of morphology formation in elastomer blends, Rub. Chem. Technol., 50, 292 (1977). 184. J.J. Elmendorp, A.K. Van Der Vegt, A study on polymer blending microrheology: Part IV. the influence of coalescence on blend morphology origination, Polym. Eng. Sci., 26, 1332–1338 (1986). ˇ y, Theory of competition between breakup and coalescence of droplets in flowing polymer 185. I. Forteln´y I, A. Zivn´ blends, Polym. Eng. Sci., 35, 1872–1877 (1995). 186. B. Majumdar, D.R. Paul, A.J. Oshinski, Evolution of morphology in compatibilized vs uncompatibilized polyamide blends, Polymer, 38, 1787–1808 (1997). 187. H. Liang, B.D. Favis, Correlation between the interfacial tension and dispersed phase morphology in interfacially modified blends of LLDPE and PVC, Macromolecules, 32, 1637–1642 (1999). 188. S.P. Lyu, F.S. Bates, C.W. Macosko, Modeling of coalescence in polymer blends, AIChE J., 48, 7–14 (2002). 189. J. Huitric, M. Moan, P.J. Carreau, N. Dufaure, Effect of reactive compatibilization on droplet coalescence in shear flow, J. Non-Newtonian Fluid Mech., 145, 139–149 (2007). 190. M. Smoluchowski, Versuch einer mathematischen theorie der koagulationskinetik kolloider l¨osungen, Z. Phys. Chem. (Zeitung F¨ur Physikalische Chemie), 92, 129–154 (1917). 191. I. Forteln´y, Z. Cern´a, J. Binko, J. Kov´ar, Anomalous dependence of the size of droplets of disperse phase on intensity of mixing, J. Appl. Polym. Sci., 48, 1731–1737 (1993). 192. R. Gonz´alez-Nu˜nez, M. Arellano, F.J. Moscoso, V.M. Gonz´alez-Romero, B.D. Favis, Determination of a limiting dispersed phase concentration for coalescence in PA6/HDPE blends under extensional flow, Polymer, 42, 5485–5489 (2001). 193. C.W. Macosko, P. Gu´egan, A.K. Khandpur, A. Nakayama, P. Marechal, T. Inoue, Compatibilizers for melt blending: premade block copolymers, Macromolecules, 29, 5590–5598 (1996). 194. S.P. Lyu, T.D. Jones, F.S. Bates, C.W. Macosko, Role of block copolymers on suppression of droplet coalescence, Macromolecules, 35, 7845–7855 (2002). 195. S.T. Milner, H.W. Xi, How copolymers promote mixing of immiscible homopolymers, J. Rheol., 40, 663–687 (1996). 196. J.W. Ha, Y. Yoon, L.G. Leal, The effect of compatibilizer on the coalescence of two drops in flow, Phys. Fluids, 15, 849–867 (2003). 197. N. Hasegawa, H. Okamoto, M. Kawasumi, A. Usuki, Preparation and mechanical properties of polystyrene-clay hybrids, J. Appl. Polym. Sci., 74, 3359–3364 (1999). 198. N. Hasegawa, H. Okamoto, M. Kato, A. Usuki, N. Sato, Nylon 6/Na–montmorillonite nanocomposites prepared by compounding nylon 6 with Na–montmorillonite slurry, Polymer, 44, 2933–2937 (2003). 199. N. Fedullo, M. Sclavons, C. Bailly, J.M. Lefebvre, J. Devaux, Nanocomposites from untreated clay: a myth? Macromol. Symp., 233, 235–245 (2006). 200. J.K. Lee, C.D. Han, Evolution of polymer blend morphology during compounding in a twin-screw extruder, Polymer, 41, 1799–1815 (2000). 201. Z.G. Cheng, Q. Wang, Morphology control of polyoxymethylene/thermoplastic polyurethane blends by adjusting their viscosity ratio, Polym. Inter., 55, 1075–1080 (2006). 202. P. P¨otschke, K. Wallheinke, H. Fritsche, H. Stutz, Morphology and properties of blends with different thermoplastic polyurethanes and polyolefines, J. Appl. Polym. Sci., 64, 749–762 (1997). 203. S. Namhata, M.J. Guest, L.M. Aerts, Blend morphology development during melt flow: Correlation of a model concept based on dynamic phase volume with practical observations, J. Appl. Polym. Sci., 71, 311–318 (1999). 204. R. Ratnagiri, C.E. Scott, Effect of viscosity variation with temperature on the compounding behavior of immiscible blends, Polym. Eng. Sci., 39, 1823–1835 (1999). 205. B.D. Favis, J.P. Chalifoux, The effect of viscosity ratio on the morphology of polypropylene/polycarbonate blends during processing, Polym. Eng. Sci., 27, 1591–1600 (1987).

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206. P.T. Hietaoja, R.M. Holsti-Miettinen, J.V. Sepp¨al¨a, O.T. Ikkala, The effect of viscosity ratio on the phase inversion of polyamide 66/polypropylene blends, J. Appl. Polym. Sci., 54, 1613–1623 (1994). 207. J. Leblanc, M. Mercier, C. Fonteix, F. Pla, Design of blends with an extremely low viscosity ratio between the dispersed and continuous phases. Dependence of the dispersed phase size on the processing parameters, Macromol. Symp., 259, 85–93 (2007). 208. M.A. Huneault, F. Mighri, G.H. Ko, F. Watanabe. Dispersion in high viscosity ratio polyolefin blends, Polym. Eng. Sci., 41, 672–683 (2001). 209. J.K. Rameshwaram, Y.S. Yang, H.S. Jeon, Structure–property relationships of nanocomposite-like polymer blends with ultrahigh viscosity ratios, Polymer, 46, 5569–5579 (2005). 210. S. Mukherjee, K. Sarkar, Effects of viscosity ratio on deformation of a viscoelastic drop in a Newtonian matrix under steady shear, J. Non-Newtonian Fluid Mech., 160, 104–112 (2009). 211. J. Kang, T.G. Smith, D.I. Bigio, Study of breakup mechanisms in cavity flow, AIChE J., 42, 649–659 (1996). 212. S. Shokoohi, A. Arefazar, A review on ternary immiscible polymer blends: morphology and effective parameters, Polym. Adv. Technol., 20, 433–447 (2009). 213. J. Reignier, B.D. Favis, M.C. Heuzey, Factors influencing encapsulation behavior in composite droplet-type polymer blends, Polymer, 44, 49–59 (2003). 214. B.K. Kim, M.S. Kim, K.J. Kim, Viscosity effect in polyolefin ternary blends and composites, J. Appl. Polym. Sci., 48, 1271–1278 (1993). 215. L.A. Utracki, M.R. Kamal, Melt rheology of polymer blends, Polym. Eng. Sci., 22, 96–114 (1982). 216. M.H. Ha, B.K. Kim, E.Y. Kim, Effects of dispersed phase composition on thermoplastic polyolefins, J. Appl. Polym. Sci., 93, 179–188 (2004). 217. M. Hemmati, H. Nazokdast, H.S. Panahi, Study on morphology of ternary polymer blends. I. Effects of melt viscosity and interfacial interaction, J. Appl. Polym. Sci., 82, 1129–1137 (2001). 218. I.H. Do, L.K. Yoon, B.K. Kim, H.M. Jeong, Effect of viscosity ratio and peroxide/coagent treatment in PP/EPR/PE ternary blends, Eur. Polym. J., 32, 1387–1393 (1996). 219. M.M. Abolhasani, A. Arefazar, M. Mozdianfard, Effect of dispersed phase composition on morphological and mechanical properties of PET/EVA/PP ternary blends, J. Polym. Sci. Part B: Polym. Phys., 48, 251–259 (2010). 220. M.H. Ha, B.K. Kim, Effects of the viscosity ratio on polyolefin ternary blends, J. Appl. Polym. Sci., 91, 4027–4036 (2004). 221. N. Nemirovski, A. Siegmann, M. Narkis, Morphology of ternary immiscible polymer blends, J. Macromol. Sci., Phys., B34, 459–475 (1995). 222. I. Luzinov, C. Pagnoulle, R. J´erˆome, Ternary polymer blend with core–shell dispersed phases: effect of the coreforming polymer on phase morphology and mechanical properties, Polymer, 41, 7099–7109 (2000). 223. G. Wildes, H. Keskkula, D.R. Paul, Coalescence in PC/SAN blends: effect of reactive compatibilization and matrix phase viscosity, Polymer, 40, 5609–5621 (1999). 224. Y. Seo, J. Kim, Structure development of TLCP ternary blends during biaxial elongational flow, Polymer, 42, 5029–5036 (2001). 225. F. Gauthier, H.L. Goldsmith, S.G. Mason, Particle motions in non-Newtonian media II. Poiseuille flow, Trans. Soc. Rheol., 15, 297–330 (1971). 226. N. Mechbal, M. Bousmina, In situ observations of unusual drop deformation and wobbling in simple shear flow, Rheol. Acta., 48, 653–663 (2009). 227. K. Verhulst, P. Moldenaers, M. Minale, Drop shape dynamics of a Newtonian drop in a non-Newtonian matrix during transient and steady shear flow, J. Rheol., 51, 261–273 (2007). 228. T.S. Valera, A.T. Morita, N.R. Demarquette, Study of morphologies of PMMA/PP/PS ternary blends, Macromolecules, 39: 2663–2675 (2006). 229. R.C. Willemse, A. Posthuma de Boer, J. van Dam, A.D. Gotsis, Co-continuous morphologies in polymer blends: the influence of the interfacial tension, Polymer, 40, 827–834 (1999). 230. J.C. Lepers, B.D. Favis, R.J. Tabar, The relative role of coalescence and interfacial tension in controlling dispersed phase size reduction during the compatibilization of polyethylene terephthalate/polypropylene blends, J. Polym. Sci. Part B: Polym. Phys., 35, 2271–2280 (1997).

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231. S. Horiuchi, N. Matchariyakul, K. Yase, T. Kitano, Morphology development through an interfacial reaction in ternary immiscible polymer blends, Macromolecules, 30, 3664–3670 (1997). 232. A.N. Wilkinson, M.L. Clemens, V.M. Harding, The effects of SEBS-g-maleic anhydride reaction on the morphology and properties of polypropylene/PA6/SEBS ternary blends, Polymer, 45, 5239–5249 (2004). 233. J. Reignier, B.D. Favis, Core–shell structure and segregation effects in composite droplet polymer blends, AIChE J., 49, 1014–1023 (2003). 234. N. Virgilio, B.D. Favis, M.F. P´epin, P. Desjardins, G. L’Esp´erance, High contrast imaging of interphases in ternary polymer blends using focused ion beam preparation and atomic force microscopy, Macromolecules, 38, 2368–2375 (2005). 235. A. Legros, P.J. Carreau, B.D. Favis, A. Michel, Morphology modification by interfacial chemical reaction in a polyester/ethylene vinyl acetate/polyethylene blend, Polymer, 38, 5085–5089 (1997). 236. J. Reignier, B.D. Favis, On the presence of a critical shell volume fraction leading to pseudo-pure droplet behavior in composite droplet polymer blends, Polymer, 44, 5061–5066 (2003). 237. M. Si, T. Araki, H. Ade, A.L.D. Kilcoyne, R. Fisher, J.C. Sokolov, M.H. Rafailovich, Compatibilizing bulk polymer blends by using organoclays, Macromolecules, 39, 4793–4801 (2006). 238. B. Crist, A.R. Nesarikar. Coarsening in polyethylene-copolymer blends, Macromolecules, 28, 890–896 (1995). 239. C. Sailer, U.A. Handge, Reactive blending of polyamide 6 and styrene-acrylonitrile copolymer: influence of blend composition and compatibilizer concentration on morphology and rheology, Macromolecules, 41, 4258–4267 (2008). 240. C. Koning, M. Van Duin, C. Pagnoulle, R. Jerome, Strategies for compatibilization of polymer blends, Prog. Polym. Sci., 23, 707–757 (1998). 241. S. Velankar, P. Van Puyvelde, J. Mewis, P. Moldenaers, Effect of compatibilization on the breakup of polymeric drops in shear flow, J. Rheol., 45, 1007–1019 (2001). 242. L. Barangi, F. Afshar Taromi, H. Nazockdast, Study on in-situ compatibilization of the PP/PA6 blends in twin screw extruders, Macromol. Symp., 274, 166–170 (2008). 243. M. Yee, P.S. Calv˜ao, N.R. Demarquette, Rheological behavior of poly(methyl methacrylate)/polystyrene (PMMA/PS) blends with the addition of PMMA-ran-PS, Rheol. Acta., 46, 653–664 (2007). 244. C.W. Macosko, H.K. Jeon, T.R. Hoye, Reactions at polymer–polymer interfaces for blend compatibilization, Prog. Polym. Sci., 30, 939–947 (2005). 245. P. Van Puyvelde, S. Velankar, P. Moldenaers, Rheology and morphology of compatibilized polymer blends, Curr. Opin. Colloid Interface Sci., 6, 457–463 (2001). 246. N. Mekhilef, B.D. Favis, P.J. Carreau, Morphological stability, interfacial tension, and dual-phase continuity in polystyrene-polyethylene blends, J. Polym. Sci. Part B: Polym. Phys., 35, 293–308 (1997). 247. P. P¨otschke, K. Wallheinke, H. Stutz, Blends of thermoplastic polyurethane and maleic-anhydride grafted polyethylene. I: morphology and mechanical properties, Polym. Eng. Sci., 39, 1035–1048 (1999). 248. D.F. Wu, Y.S. Zhang, L.J. Yuan, M. Zhang, W.D. Zhou, Viscoelastic interfacial properties of compatibilized poly(ε-caprolactone)/polylactide blend, J. Polym. Sci. Part B: Polym. Phys., 48, 756–765 (2010). 249. P. Cigana, B.D. Favis, R. J´erˆome, Diblock copolymers as emulsifying agents in polymer blends: influence of molecular weight, architecture, and chemical composition, J. Polym. Sci. Part B: Polym. Phys., 34, 1691–1700 (1996). 250. S. Filippi, L. Minkova, N. Dintcheva, P. Narducci, P. Magagnini, Comparative study of different maleic anhydride grafted compatibilizer precursors towards LDPE/PA6 blends: morphology and mechanical properties, Polymer, 46, 8054–8061 (2005). 251. H. Cartier, G.H. Hu, Morphology development of in situ compatibilized semicrystalline polymer blends in a corotating twin-screw extruder, Polym. Eng. Sci., 39, 996–1013 (1999). 252. P. Martin, J. Devaux, R. Legras, L. Leemans, M. van Gurp, M. van Duin, Complex processing-morphology interrelationships during the reactive compatibilization of blends of poly(butylene terephthalate) with epoxide-containing rubber, J. Appl. Polym. Sci., 91, 703–713 (2004). 253. F. Abbassi-Sourki, M.A. Huneault, M. Bousmina, Effect of compatibilization on the deformation and breakup of drops in step-wise increasing shear flow, Polymer, 50, 645–653 (2009).

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254. L. Levitt, C.W. Macosko, Shearing of polymer drops with interface modification, Macromolecules, 32, 6270–6277 (1999). 255. P. Van. Puyvelde, S. Velankar, J. Mewis, P. Moldenaers, Effect of marangoni stresses on the deformation and coalescence in compatibilized immiscible polymer blends, Polym. Eng. Sci., 42, 1956–1964 (2002). 256. H.K. Jeon, C.W. Macosko, Visualization of block copolymer distribution on a sheared drop, Polymer, 44, 5381–5386 (2003). 257. S. Cimmino, F. Coppola, L. D’Orazio, R. Greco, G. Maglio, M. Malinconico, C. Mancarella, E. Martuscelli, G. Ragosta, Ternary nylon-6/rubber/modified rubber blends: effect of the mixing procedure on morphology, mechanical and impact properties, Polymer, 27, 1874–1884 (1986). 258. M.B. Coltelli, E. Passaglia, F. Ciardelli, One-step functionalization and reactive blending of polyolefin/polyamide mixtures (EPM/PA6), Polymer, 47, 85–97 (2006). 259. D. Tomova, H.J. Radusch, Morphology and properties of ternary polyamide 6/polyamide 66/elastomer blends, Polym, Adv. Technol., 14, 19–26 (2003). 260. C.A. de Freitas, T.S. Valera, A.M.C. Souza, N.R. Demarquette, Morphology of compatibilized ternary blends, Macromol. Symp., 247, 260–270 (2007). 261. N. Virgilio, P. Sarazin, B.D. Favis, Towards ultraporous poly(L-lactide) scaffolds from quaternary immiscible polymer blends, Biomaterials, 31, 5719–5728 (2010). 262. M. Zanetti, L. Costa, Preparation and combustion behavior of polymer/layered silicate nanocomposites based upon PE and EVA, Polymer, 45, 4367–4373 (2004). 263. S. Hotta, D.R. Paul, Nanocomposites formed from linear low density polyethylene and organoclays, Polymer, 45, 7639–7654 (2004). 264. K.H. Wang, M.H. Choi, C.M. Koo, Y.S. Choi, I.J. Chung, Synthesis and characterization of maleated polyethylene/clay nanocomposites, Polymer, 42, 9819–9826 (2001). 265. A. Usuki, M. Kato, A. Okada, T. Kurauchi, Synthesis of polypropylene–clay hybrid, J. Appl. Polym. Sci., 63, 137–139 (1997). 266. N. Hasegawa, M. Kawasumi, M. Kato, A. Usuki, A. Okada, Preparation and mechanical properties of polypropyleneclay hybrids using a maleic anhydride-modified polypropylene oligomer, J. Appl. Polym. Sci., 67, 87–92 (1998). 267. N. Hasegawa, A. Usuki, Silicate layer exfoliation in polyolefin/clay nanocomposites based on maleic anhydride modified polyolefins and organophilic clay, J. Appl. Polym. Sci., 93, 464–470 (2004). 268. G. Galgali, S. Agarwal, A.K. Lele, Effect of clay orientation on the tensile modulus of polypropylene-nanoclay composites, Polymer, 45, 6059–6069 (2004). 269. P.H. Nam, P. Maiti, M. Okamoto, T. Kotaka, N. Hasegawa, A. Usuki, A hierarchical structure and properties of intercalated polypropylene/clay nanocomposites, Polymer, 42, 9633–9640 (2001). 270. M.L. L´opez-Quintanilla, S. S´anchez-Vald´es, L.F. Ramos de Valle, F.J. Medell´ın-Rodr´ıguez, Effect of some compatibilizing agents on clay dispersion of polypropylene-clay nanocomposites, J. Appl. Polym. Sci., 100, 4748–4756 (2006). 271. M. Kato, H. Okamoto, N. Hasegawa, A. Tsukigase, A. Usuki, Preparation and properties of polyethylene-clay hybrids, Polym. Eng. Sci., 43, 1312–1316 (2003). 272. J. Zhang, C.A. Wilkie, Preparation and flammability properties of polyethylene-clay nanocomposites, Polym. Degrad. Stabil., 80, 163–169 (2003). 273. F.C. Chiu, H.Z. Yen, C.E. Lee, Characterization of PP/HDPE blend-based nanocomposites using different maleated polyolefins as compatibilizers, Polym. Test., 29, 397–406 (2010). 274. D. Wang, C.A. Wilkie, In-situ reactive blending to prepare polystyrene-clay and polypropylene-clay nanocomposites, Polym. Degrad. Stabil., 80, 171–182 (2003). 275. B.D. Favis, J.P. Chalifoux, Influence of composition on the morphology of polypropylene/polycarbonate blends, Polymer, 29, 1761–1767 (1988). 276. B.D. Favis, D. Therrien, Factors influencing structure formation and phase size in an immiscible polymer blend of polycarbonate and polypropylene prepared by twin-screw extrusion, Polymer, 32, 1474–1481 (1991). 277. O. Franzheim, T. Rische, M. Stephan, W. J. Macknight, Blending of immiscible polymers in a mixing zone of a twin screw extruder-effects of compatibilization, Polym. Eng. Sci., 40, 1143–1156 (2000).

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278. Y.G. Cho, M.R. Kamal, Effect of the dispersed phase fraction on particle size in blends with high viscosity ratio, Polym. Eng. Sci., 42, 2005–2015 (2002). 279. G. Serpe, J. Jarrin, F. Dawans, Morphology-processing relationships in polyethylene-polyamide blends, Polym. Eng. Sci., 30, 553–565 (1990). 280. M. Hemmati, H. Nazokdast, H.S. Panahi, Study on morphology of ternary polymer blends. II. effect of composition, J. Appl. Polym. Sci., 82, 1138–1146 (2001). 281. Y. Han, Z.X. Tai, C. Zhou, M.Y. Zhang, H.X. Zhang, F.Q. Liu, Influence of blend composition on the mechanical properties and morphology of PC/ASA/SAN ternary blends, Polym. Bull., 62, 855–866 (2009). 282. J. Zhang, S. Ravati, N. Virgilio, B.D. Favis, Ultralow percolation thresholds in ternary cocontinuous polymer blends, Macromolecules, 40, 8817–8820 (2007). 283. M. Mehrabzadeh, M.R. Kamal, Polymer-clay nanocomposites based on blends of polyamide-6 and polyethylene, Can. J. Chem. Eng., 80, 1083–1092 (2002). 284. G. Filippone, N.Tz. Dintcheva, F.P. La Mantia, D. Acierno, Selective localization of organoclay and effects on the morphology and mechanical properties of LDPE/PA11 blends with distributed and co-continuous morphology, J. Polym. Sci. Part B: Polym. Phys., 48, 600–609 (2010). 285. J.S. Hong, H. Namkung, K.H. Ahn, S.J. Lee, C. Kim, The role of organically modified layered silicate in the breakup and coalescence of droplets in PBT/PE blends, Polymer, 47, 3967–3975 (2006). 286. H.S. Lee, P.D. Fasulo, W.R. Rodgers, and D.R. Paul, TPO based nanocomposites. Part 1. morphology and mechanical properties, Polymer, 46, 11673–11689 (2005). 287. S. Mehta, F.M. Mirabella, K. Rufener, A. Bafna, Thermoplastic olefin/clay nanocomposites: morphology and mechanical properties, J. Appl. Polym. Sci., 92, 928–936 (2004). 288. W. Li, J. Karger-Kocsis, R. Thomann, Compatibilization effect of TiO2 nanoparticles on the phase structure of PET/PP/TiO2 nanocomposites, J. Polym. Sci. Part B: Polym. Phys., 47, 1616–1624 (2009). 289. A. Bandyopadhyay, V. Thakur, S. Pradhan, A.K. Bhowmick, Nanoclay distribution and its influence on the mechanical properties of rubber blends, J. Appl. Polym. Sci., 115, 1237–1246 (2010). 290. D.P. Dharaiya, S.C. Jana, Nanoclay-induced morphology development in chaotic mixing of immiscible polymers, J. Polym. Sci. Part B: Polym. Phys., 43, 3638–3651 (2005). 291. C.G. Martins, N.M. Larocca, D.R. Paul, L.A. Pessan, Nanocomposites formed from polypropylene/EVA blends, Polymer, 50, 1743–1754 (2009). 292. L.Y. Zhang, C.Y. Wan, Y. Zhang, Morphology and electrical properties of polyamide 6/polypropylene/multi-walled carbon nanotubes composites, Compos. Sci. Technol., 69, 2212–2217 (2009). 293. A.C. Baudouin, J. Devaux, C. Bailly, Localization of carbon nanotubes at the interface in blends of polyamide and ethylene–acrylate copolymer, Polymer, 51, 1341–1354 (2010). 294. Y. Yoo, L. Cui, P.J. Yoon, D.R. Paul, Morphology and mechanical properties of rubber toughened amorphous polyamide/MMT nanocomposites, Macromolecules, 43, 615–624 (2010). 295. R. Gallego, D. Garc´ıa-L´opez, J.C. Merino, J.M. Pastor, The effect of montmorillonite and compatibilizer quantities on stiffness and toughness of polyamide nanoblends, Polym. Inter., 59, 472–478 (2010). 296. H.X. Huang, Y.F. Huang, S.L. Yang, Developing laminar morphology in higher viscosity ratio polyblends by controlling flow fields in a single screw extruder, Polym. Inter., 54, 65–69 (2005). 297. H.X. Huang, HDPE/PA-6 blends: Influence of screw shear intensity and HDPE melt viscosity on phase morphology development, J. Mater. Sci., 40, 1777–1779 (2005). 298. M. Heindl, M.K. Sommer, H. M¨unstedt, Morphology development in polystyrene/polyethylene blends during uniaxial elongational flow, Rheol. Acta., 44, 55–70 (2004). 299. D.A. Zumbrunnen, C. Chhibber, Morphology development in polymer blends produced by chaotic mixing at various compositions, Polymer, 43, 3267–3277 (2002). 300. D. A. Zumbrunnen, S. Inamdar, O. Kwon, P. Verma, Chaotic advection as a means to develop nanoscale structures in viscous melts, Nano. Lett., 2, 1143–1148 (2002). 301. C. Lacroix, M. Grmela, P.J. Carreau, Morphological evolution of immiscible polymer blends in simple shear and elongational flows, J. Non-Newtonian Fluid Mech., 86, 37–59 (1999). 302. H. Gramespacher, J. Meissner, Melt elongation and recovery of polymer blends, morphology and influence of interfacial tension, J. Rheol., 41, 27–44 (1997).

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303. F. Oosterlinck, I. Vinckier, M. Mours, H.M. Laun, P. Moldenaers, Morphology development of a PS/PMMA polymer blend during flow in dies, Rheol. Acta., 44, 631–643 (2005). 304. F. Oosterlinck, M. Mours, H.M. Laun, P. Moldenaers, Morphology development of a polystyrene/ polymethylmethacrylate blend during start-up of uniaxial elongational flow, J. Rheol., 49, 897–918 (2005). 305. Z. Star´y, H. M¨unstedt, Morphology development in PS/LLDPE blend during and after elongational deformation, J. Polym. Sci. Part B: Polym. Phys., 46, 16–27 (2008). 306. H.X. Huang, G. Jiang, S.Q. Mao, Microstructure and on-line shear viscosity of PP/nano-CaCO3 composites prepared by twin-screw extruder, J. Mater. Sci., 41, 4985–4988 (2006). 307. G. Jiang, H.X. Huang, Online shear viscosity and microstructure of PP/nano-CaCO3 composites produced by different mixing types, J. Mater. Sci., 43, 5305–5312 (2008). 308. G. Jiang, H.X. Huang, Microstructure and rheologic development of polypropylene/nano-CaCO3 composites along twin-screw extruder, J. Appl. Polym. Sci., 114, 1687–1693 (2009). 309. R.K. Gupta, V. Pasanovic-Zujo, S.N. Bhattacharya, Shear and extensional rheology of EVA/layered silicatenanocomposites, J. Non-Newtonian Fluid Mech., 128, 116–125 (2005). 310. M. Tokihisa, K. Yakemoto, T. Sakai, L.A. Utracki, M. Sepehr, J. Li, Y. Simard, Extensional flow mixer for polymer nanocomposites, Polym. Eng. Sci., 46, 1040–1050 (2006). 311. L.A. Utracki, M. Sepehr, J. Li, Melt compounding of polymeric nanocomposites, Inter. Polym. Process., 21, 3–16 (2006). 312. F.P. La Mantia, N.T. Dintcheva, R. Scaffaro, R. Marino, Morphology and properties of polyethylene/clay nanocomposite drawn fibers, Macromol. Mater. Eng., 293, 83–91 (2008). 313. F.P. La. Mantia, R. Marino, N.T. Dintcheva, Morphology modification of polyethylene/clay nanocomposite samples under convergent flow, Macromol. Mater. Eng., 294, 575–581 (2009). 314. E. Garofalo, G.M. Russo, P. Scarfato, L. Incarnato, Nanostructural modifications of polyamide/MMT hybrids under isothermal and nonisothermal elongational flow, J. Polym. Sci. Part B: Polym. Phys., 47, 981–993 (2009). 315. U.A. Handge, P. P¨otschke, Deformation and orientation during shear and elongation of a polycarbonate/carbon nanotubes composite in the melt, Rheol. Acta., 46, 889–898 (2007). 316. X Q. Nguyen, L. A. Utracki, Extensional Flow Mixer, US Patent, 5,451, 106 (1995). 317. A. Luciani, L.A. Utracki, The Extensional flow mixer, EFM, Inter. Polym. Process., 11, 299–309 (1996). 318. C. Testa, I. Sigillo, N. Grizzuti, Morphology evolution of immiscible polymer blends in complex flow fields, Polymer, 42, 5651–5659 (2001). 319. P.L. Maffettone, M. Minale, Equation of change for ellipsoidal drops in viscous flow, J. Non-Newtonian Fluid. Mech., 78, 227–241 (1998). 320. N. Nassar, L.A. Utracki, M.R. Kamal, Melt intercalation in montmorillonite/polystyrene nanocomposites, Inter. Polym. Process., 20, 423–431 (2005). 321. I. Delaby, B. Ernst, R. Muller, Drop deformation during elongational flow in blends of viscoelastic fluids: small deformation theory and comparison with experimental results, Rheol. Acta., 34, 525–533 (1995). 322. I. Delaby, B. Ernst, D. Froelich, R. Muller, Droplet deformation in immiscible polymer blends during transient uniaxial elongational flow, Polym. Eng. Sci., 36, 1627–1635 (1996). 323. E.M. Toose, D. van den Ende, B.J. Geurts, J.G.M. Kuerten, Axisymmetric non-Newtonian drops treated with a boundary integral method, J. Eng. Math., 30, 131–150 (1996). 324. R. W. Hooper, V.F. de Almeida, C.W. Macosko, J.J. Derby, Transient polymeric drop extension and retraction in uniaxial extensional flows, J. Non-Newtonian Fluid Mech., 98, 141–168 (2001). 325. D.A. Zumbrunnen, R. Subrahmanian, B. Kulshreshtha, C. Mahesha, Smart blending technology enabled by chaotic advection, Adv. Polym. Technol., 25, 152–169 (2006). 326. J.M. Ottino, The Kinematics of Mixing: Stretching, Chaos, and Transport, Cambridge University Press, Cambridge, 1990. 327. S.J. Kim, T.H. Kwon, Enhancement of mixing performance of single-screw extrusion processes via chaotic flows: Part II. numerical study, Adv. Polym. Technol., 15, 55–69 (1996). 328. H. Cheng, I. Manas-Zloczower, Chaotic features of flow in polymer processing equipment-relevance to distributive mixing, Inter. Polym. Process., 12, 83–91 (1997).

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329. J. F. Muzzio, M. Tjahjadi, J.M. Ottino, Self-similar drop-size distributions produced by breakup in chaotic flows, Phys. Rev. Lett., 67, 54–57 (1991). 330. R. I. Danescu, D.A. Zumbrunnen, Creation of conducting networks among particles in polymer melts by chaotic mixing, J. Thermoplast. Comp. Mater., 11, 299–320 (1998). 331. Y.H. Liu, D.A. Zumbrunnen, Toughness enhancement in polymer blends due to the in-situ formation by chaotic mixing of fine-scale extended structures, J. Mater. Sci., 34, 1921–1931 (1999). 332. O. Kwon, D.A. Zumbrunnen, Progressive morphology development to produce multilayer films and interpenetrating blends by chaotic mixing, J. Appl. Polym. Sci., 82, 1569–1579 (2001). 333. S.C. Jana, M. San, Effects of viscosity ratio and composition on development of morphology in chaotic mixing of polymers, Polymer, 45, 1665–1678 (2004). 334. M. San, S.C. Jana, A study on the effects of chaotic mixer design and operating conditions on morphology development in immiscible polymer systems, Polym. Eng. Sci., 44, 407–422 (2004). 335. A. Dhoble, B. Kulshreshtha, S. Ramaswami, D.A. Zumbrunnen, Mechanical properties of PP-LDPE blends with novel morphologies produced with a continuous chaotic advection blender, Polymer, 46, 2244–2256 (2005). 336. H.X. Huang, Y.C. Peng, Theoretical modeling of dispersive melting mechanism of polymers in an extruder, Adv. Polym. Technol., 12, 343–352 (1993). 337. S.J. Kim, T.H. Kwon, Enhancement of mixing performance of single-screw extrusion processes via chaotic flows: Part I. basic concepts and experimental study, Adv. Polym. Technol., 15, 41–54 (1996). 338. R.K. Thakur, Ch. Vial, K.D.P. Nigam, E.B. Nauman, G. Djelveh, Static mixers in the process industries – a review, Trans IChemE, Part A, Chem. Eng. Res. Des., 81, 787–826 (2003). 339. D.A. Zumbrunnen, S. Inamdar. Novel sub-micron highly multi-layered polymer films formed by continuous flow chaotic mixing, Chem. Eng. Sci., 56, 3893–3897 (2001). 340. X. Jin, D.A. Zumbrunnen, Morphology development by chaotic advection in PA-EVOH blends with polymer component interactions, SPE ANTEC Tech. Papers, 1470–1474 (2008). 341. O. Kwon, D.A. Zumbrunnen, Production of barrier films by chaotic mixing of plastics, Polym. Eng. Sci., 43, 1443–1459 (2003). 342. D.P. Dharaiya, S.C. Jana, S.F. Lyuksyutov, Production of electrically conductive networks in immiscible polymer blends by chaotic mixing, Polym. Eng. Sci., 46, 19–28 (2006). 343. G.A. Jimenez, S.C. Jana, Electrically conductive polymer nanocomposites of polymethylmethacrylate and carbon nanofibers prepared by chaotic mixing, Compos. Part A: Appl. Sci. Manufacturing, 38, 983–993 (2007). 344. H.X. Huang, Y.F. Huang, C.Y. Wang, Y.H. Zhang. Microstructure changes of a polypropylene/montmorillonite nanocomposite, J. Mater. Sci. Lett., 22, 1547–1549 (2003). 345. H.X. Huang, C.Y. Wang, Y.F. Huang, G. Jiang, Y.H. Zhang, Effect of mixing element in a single screw extruder on the microstructure of polypropylene/montmorillonite nanocomposites, SPE ANTEC Tech. Papers, 2223–2227 (2003). 346. D. Bourry, B.D. Favis, Morphology development in a polyethylene/polystyrene binary blend during twin-screw extrusion, Polymer, 39, 1851–1856 (1998). 347. M.H. Ha, M.S. Kim, B.K. Kim, W. Kim, M.C. Lee, H.D. Kim, Effects of the blending sequence in polyolefin ternary blends, J. Appl. Polym. Sci., 92, 804–811 (2004). 348. J.J. Huang, H. Keskkula, D.R. Paul, Elastomer particle morphology in ternary blends of maleated and non-maleated ethylene-based elastomers with polyamides: role of elastomer phase miscibility, Polymer, 47, 624–638 (2006). 349. R. Gallego, D. Garc´ıa-L´opez, S. L´opez-Quintana, I. Gobernado-Mitre, J.C. Merino, J.M. Pastor, Influence of blending sequence on micro- and macrostructure of PA6/mEPDM/EPDMgMA blends reinforced with organoclay, J. Appl. Polym. Sci., 109, 1556–1563 (2008). 350. I. Gonz´alez, J.I. Eguiaz´abal, J. Naz´abal, Effects of the processing sequence and critical interparticle distance in PA6-clay/mSEBS nanocomposites, Eur. Polym. J., 44, 287–299 (2008). 351. S. Filippi, N.T. Dintcheva, R. Scaffaro, F.P. La. Mantia, G. Polacco, P. Magagnini, Effects of organoclay on morphology and properties of nanocomposites based on LDPE/PA-6 blends without and with SEBS-g-MA compatibilizer, Polym. Eng. Sci., 49, 1187–1197 (2009). 352. Y. Wang, F.B. Chen, K.C. Wu, Twin-screw extrusion compounding of polypropylene/organoclay nanocomposites modified by maleated polypropylenes, J. Appl. Polym. Sci., 93, 100–112 (2004).

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353. X.C. Li, H.M. Park, J.O. Lee, C.S. Ha, Effect of blending sequence on the microstructure and properties of PBT/EVA-g-MAH/organoclay ternary nanocomposites, Polym. Eng. Sci., 42, 2156–2164 (2002). 354. D. Wu, C. Zhou, W. Yu, F. Xie, Effect of blending sequence on the morphologies of poly(butylene terephthalate)/epoxy/clay nanocomposites by a rheological approach, J. Appl. Polym. Sci., 99, 340–346 (2006). 355. B. Akbari, R. Bagheri, Influence of PP-g-MA on morphology, mechanical properties and deformation mechanism of copolypropylene/clay nanocomposite, J. Appl. Polym. Sci., 114, 3751–3759 (2009). 356. K. Prakashan, A. K. Gupta, S. N. Maiti, Effect of the mixing sequence on the morphology and properties of a polypropylene/polydimethylsiloxane/nano-SiO2 Ternary Composite, J. Appl. Polym. Sci., 110, 1457–1468 (2008). 357. J.J. Huang, H. Keskkula, D.R. Paul, Rubber toughening of an amorphous polyamide by functionalized SEBS copolymers: morphology and Izod impact behavior, Polymer, 45, 4203–4215 (2004). 358. V. Bordereau, M. Carrega, Z. H. Shi, L. A. Utracki, P. Sammut, Development of polymer blend morphology during compounding in a twin-screw extruder: Part III: experimental procedure and preliminary results, Polym. Eng. Sci., 32, 1846–1856 (1992). 359. M. Modesti, A. Lorenzetti, D. Bon, S. Besco, Effect of processing conditions on morphology and mechanical properties of compatibilized polypropylene nanocomposites, Polymer, 46, 10237–10245 (2005). 360. W. Lertwimolnun, B. Vergnes, Influence of compatibilizer and processing conditions on the dispersion of nanoclay in a polypropylene matrix, Polymer, 46, 3462–3471 (2005). 361. W. Lertwimolnun, B. Vergnes, Effect of processing conditions on the formation of polypropylene/organoclay nanocomposites in a twin screw extruder, Polym. Eng. Sci., 46, 314–323 (2006). 362. P. Peltola, E. V¨alipakka, J. Vuorinen, S. Syrj¨al¨a, and K. Hanhi, Effect of rotational speed of twin screw extruder on the microstructure and rheological and mechanical properties of nanoclay-reinforced polypropylene nanocomposites, Polym. Eng. Sci., 46, 995–1000 (2006). 363. M.D Read., L. Liu., J.D. Harris, R.R. Samson, The development of structure in a melt blended polypropylene organoclay nanocomposite, SPE ANTEC Tech. Papers, 1882–1886 (2004). 364. T. McNally, W.R. Murphy, C.Y. Lew, R.J. Turner, G.P. Brennan, Polyamide-12 layered silicate nanocomposites by melt compounding, Polymer, 44, 2761–2772 (2003). 365. Z.M. Liang, J. Yin, Poly(etherimide)/montmorillonite nanocomposites prepared by melt intercalation, J. Appl. Polym. Sci., 90, 1857–1863 (2003). 366. E.K. Hobbie, H. Wang, H. Kim, S. Lin-Gibson, E.A. Grulke, Orientation of carbon nanotubes in a sheared polymer melt, Phys. Fluids., 15, 1196–1202 (2003).

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7 Mechanical and Viscoelastic Characterization of Multiphase Polymer Systems Poornima Vijayan P. School of Chemical Sciences, Mahatma Gandhi University, Kerala, India

Siby Varghese Rubber Research Institute of India, Kottayam, Kerala, India

Sabu Thomas Centre for Nanoscience and Nanotechnology, Mahatma Gandhi University, Kerala, India

7.1 Introduction Multiphase polymer systems, which form a fast-growing development in the area of polymer science and technology, include blends, composites, alloys, interpenetrating polymer networks (IPNs) and gels. Polymer blends are mixtures of two or more polymers and/or copolymers in which the minor component contributes at least 2 wt%. Polymer blends have gained significant growth in the last two decades and they constitute ca. 36 wt% of total polymer consumption. Current worldwide market volume for polymer blends and alloys is estimated to be more than 700,000 metric tons/year with an average growth rate of 6 to 7 % [1–6]. Figure 7.1 illustrates the price/performance of commercial polymer blends. Commercial polymer blends have existed for decades; however, the concepts of miscibility, phase behavior and the basic nature of polymer blends were not well understood. It is perhaps that the successful commercialization of miscible blends of poly (2,6-dimethyl-1,4 phenylene ether) and polystyrene prompted the interest in polymer blends. The inspiration for polymer scientists and industrialists to focus on polymer blends rather than synthesizing new materials arose from cheaper development costs, maximum diversification and

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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SPECIALITY POLYMERS

PRICE

PEEK, LCP, PES, PSU, PAR, PEI, PPS

PERFORMANCE

ENGINEERING THERMOPLASTICS PA, PBS, PET, PC, POM, PPO, ABS, SMA, Acrylics

COMMODITY THERMOPLASTICS HDPE, LDPE, LLDPE, PP, PS, HIPS, PVC

CONSUMPTION

Figure 7.1 The price/quality profile of commercial polymer blends. Reprinted from [1]. Copyright (2002) with permission from Springer.

increased use of existing polymers. Additional benefits are: (1) improved specific properties; (2) improved means for the industrial recycling; (3) plant flexibility and productivity; and (4) diversified applications of polymer blends which include automotive, electronics and electrical, building and constructions, medical and packaging, etc. Recently, nanostructured polymer blend systems have become increasingly important. In nanoblends, the scale of dispersion of one polymer into another is in general below 100 nm.Nanostructured polymer blends very often exhibit unique properties that are directly attributed to the presence of structural entities having dimension in the nanometer range [7]. On the other hand, interpenetrating polymer networks (IPNs) are a blend of two or more polymers in a network form, at least one of which is synthesized and/or crosslinked in the immediate presence of the other(s). When one of the polymers is crosslinked, the product is called semi IPNs. An IPN can be distinguished from polymer blends, blocks, or grafts in two ways: (1) an IPN swells, but does not dissolve in solvents, and (2) creep and flow are suppressed [1]. The IPN is held together exclusively by mutually entangled bands of nanometer scale dimension. This kind of ‘supramolecular’structure favors shear yielding, which is the major energy dissipation mechanism in thermosets [8]. Existence of the physical interlocking (interlocked macrocycles of the two component networks) had been the major factor in determining the characteristic properties [9]. IPN structuring offers unique possibilities for composites with respect to fiber-to-matrix adhesion [8]. Another important and widely-used multiphase polymer system is polymer gels. A polymer gel consists of an elastic crosslinked network and a fluid filling the interstitial spaces of the network. The network of long polymer molecules holds the liquid in place and so gives the gel what solidity it has [10]. Polymer gels can be easily deformed by external stimuli, and generate force or execute work on the external environment. If such responses can be translated from the microscopic level to a macroscopic scale, a conversion of chemical free energy into mechanical work should be realized. The ability of polymer gels to undergo substantial swelling and collapsing as a function of their environment is one of the most remarkable properties of these materials. The phenomenon of gel volume transitions, which can be induced by temperature, pH, or ionic strength, has prompted researchers to investigate gels as potential actuators, sensors controllable membranes for separations, and modulators for delivery of drugs [11].

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Polymer science uses the concept of a composite as a material formed by combining different phases. In conventional polymer composites, many types of inorganic fillers with dimensions in the micrometer range, e.g., calcium carbonate, glass beads, etc. have been used extensively to enhance the mechanical properties of polymers. A nanocomposite is defined as a composite material where at least one of the dimensions of one of its constituents is in the nanometer scale. How nanoparticles affect the properties of nanocomposite materials is of concern to engineers and scientists who are interested in the applications of nanocomposites [12]. Final properties of the multiphase polymer system are strongly controlled by various factors such as morphology, dispersion, mixing conditions, filler content, etc. This chapter discusses the various parameters that affect the mechanical properties of the above-mentioned multiphase polymer systems separately. Various models have been used to predict the mechanical and viscoelastic properties.

7.2 Polymer Blends 7.2.1

Ultimate Mechanical Properties and Modeling

Major parameters which control the mechanical properties of the polymer blends are: morphology, mixing conditions, blend composition, crosslinking, compatibilization and filler addition, etc. 7.2.1.1

Influence of Blend Morphology on Mechanical Properties

The final properties of polymer blends are greatly dependent on the morphology of the systems. Morphology of a polymer blend indicates the size, shape and spatial distribution of the component phases with respect to each other. Since it is well established that most of the properties of the polymer blend are strongly influenced by the type and fineness of the phase structure, the study of the control of the morphology of the polymer blend has emerged as an area of continuous interest to polymer material scientists in the last few decades [13, 14]. For a given blend, various types of useful morphologies (Figure 7.2) for different end properties such as high strength and toughness, toughness coupled with stiffness, good barrier properties, and high flow can be obtained [14]. From the point of view of broader classification, multiphase polymer blends may be divided into two major categories: 1. Blends with a discrete phase structure (i.e. droplets in matrix). 2. Blends with a bicontinuous phase structure (i.e. co-continuous). Blends with a discrete phase structure are most common in which droplets of the minor phase dispersed in a matrix of the major phase. These types of blends are often used in rubber modification of brittle polymers. The minor phase can also be dispersed as fibers, for example in self-reinforcing polymer blends. In these kinds of blends, the properties are mainly improved in the direction of the fibers. In co-continuous morphologies, the interesting feature is that both components, in all directions, can fully contribute to the properties of the blend. Attempts have been made by Thomas and co-workers to correlate the blend morphology with the observed mechanical properties [15, 16]. Varghese et al. studied the correlation between morphology and mechanical properties of nitrile rubber/ethylene-vinyl acetate copolymer (NBR/EVA) [15]. Morphology of the NBR/EVA blends indicated a two-phase structure in which the minor phase is dispersed as domains in the major continuous phase. However, between 40 and 50 wt% of NBR content both NBR and EVA exist as continuous phases and generate a co-continuous morphology. The stress-strain curves of EVA-rich blends have a similar

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Drops

(toughness, surface modification)

1 μm

Double emulsion

Laminar

(toughness and stiffness)

(barrier)

Fibers

Cocontinuous

(strength, thermal expansion)

(high flow, electrical conductivity, toughness, stiffness)

Ordered microphases

10 μm

Figure 7.2 Schematic of useful morphologies of polymer blends. Reprinted from [8]. Copyright (2006) Taylor and Francis Group.

behavior as that of pure EVA, NBR rich blends as that of typical unfilled rubber and the 50–50 blend with a co-continuous morphology showed an intermediate stress–strain behavior. Tensile strength and elongation at break increase with the increase in EVA content. The increase was found to be sharper when the EVA content was more than 40% where it formed a continuous phase. Kumar et al. found that the morphology of the blend has a strong influence on the mechanical properties of nylon/NBR blend [16]. The mechanical properties were found to increase rapidly beyond 40 wt % of nylon (Figure 7.3). This abrupt rise in mechanical properties is associated with the fully co-continuous nature of the nylon matrix (Figure 7.4). Veenstra et al. measured and compared the mechanical properties of polymer blends with co-continuous phase morphologies to the properties of blends of the same polymers with a droplet-matrix morphology [17]. In their study, PS/poly(ether-ester) and polypropylene/styrene-butylene-styrene (PP/SEBS) copolymer blends were prepared with both morphologies (dispersed-matrix and co-continuous). Elastic moduli of the cocontinuous blends were significantly higher than the moduli of the dispersed blends. However, no significant difference in tensile strength was found when the co-continuous blends were compared to blends with droplet-matrix morphology.

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MAXIMUM TENSILE STRENGTH, MPa

50

40

30

20

10

0

0

20

40

60

80

100

WEIGHT % OF NYLON

Figure 7.3 Effect of nylon/NBR blend composition on maximum tensile strength. Reprinted from [16]. Copyright (1996) with permission from John Wiley & Sons.

(a)

(b)

20.3 μm 7.1 μm

(c)

7.7 μm

Figure 7.4 SEM images of the blend morphology of (a) N30 showing the dispersed nylon particles in the continuous NBR matrix. (b) N50 showing a co-continuous morphology. (c) N70 here NBR particles are dispersed in the continuous nylon matrix. Reprinted from [17]. Copyright (2000) with permission from Elsevier.

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(a) 2000

3000

1000 1000 Ey (MPa)

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30 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 Volume fraction PS

100

10

1 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 Volume fraction PP

Figure 7.5 (a) Young’s moduli for the PS/poly(ether–ester) systems Ia () and Ib (•) as a function of volume fraction PS; (b) Young’s moduli for the PP/SEBS systems IIa () and IIb (•) as a function of volume fraction PP (when the error bars are not visible they are smaller than the marker) [17].

In Figure 7.5a, the Young’s moduli for two blend systems of PS/poly(ether-ester), Ia (the blends processed at 230◦ C) and Ib (the blends processed at 200◦ C), are plotted as a function of PS. Morphology of the Ia showed co-continuity over a small composition range (50 to 60 vol % PS), and the morphology of the Ib blends showed a broad range of co-continuity (30–60vol %PS) (Figure 7.6). All other compositions of Ia and Ib blends showed droplet-matrix morphologies. Moduli show a high increase when PS becomes continuous throughout the sample. It is important to note that when the moduli of the co-continuous blends are compared to that of the dispersed morphologies (with the same volume fractions), it becomes evident that at low volume fractions of PS, the co-continuous blends show higher values for the Young’s modulus than the dispersed blends. It is very clear that PS contributes more to the modulus of the blend when it is continuous than when it is dispersed in the poly(ether-ester) matrix. This is associated with the fact that in co-continuous morphologies both phases take part in the load-bearing process. At higher volume fractions of PS, the difference in modulus related to the morphology diminishes. Similarly, in Figure 7.5b, the Young’s moduli for two blend systems of PP/SEBS blends, IIa (the blends processed at 250◦ C), and IIb (the blends processed at 190◦ C), are plotted as a function of the volume fraction of PP. The morphology of the IIa blends showed continuity over a small composition range (50 to 60 vol% PP), and the morphology of the IIb blends showed a broad range of co-continuity (40 to 80 vol % PP). From this, it is evident that co-continuous blends show a much higher value for Young’s modulus than the dispersed blends.

7.2.1.2

Influence of Mixing Conditions on Mechanical Properties

The mechanical properties of polymer blends strongly depend on processing conditions (temperature, time, intensity and type of mixing and nature of flow). Mechanical properties of natural rubber/poly(methyl methacrylate) blends were investigated as a function of mixing conditions by Oommen and Thomas [18]. A comparison of tensile strength and tear strength of the solution-cast and melt-mixed samples was carried out. Even though both systems exhibit considerable improvement in mechanical strength, the melt-mixed

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Figure 7.6 SEM micrographs of PS/poly(ether–ester) blends of systems Ia (a and b) and Ib (c and d) with 30 and 40 vol.% PS [17].

samples were found to show lower strength values as compared to respective solution-cast samples (Table 7.1). Melt- mixed blends exhibit lower strength both in the compatibilized and uncompatibilized states compared to the solution-cast system. The improved strength of the solution-cast blend is because molecular-level mixing is achieved during solution mixing and this leads to improved adhesion between the phases. The high shearing action (80 rpm) and temperature (145◦ C) employed during the preparation of the blend by the melt-mixing technique might have caused degradation of NR and PMMA, resulting in substantial reduction in strength. 7.2.1.3

Blend Composition

Blend composition of the multiphase polymer system is one of the important factors which controls the morphology of the system and hence the mechanical properties. The mechanical properties of PLA–PCL blends can be tuned through the blend composition. Blends of biodegradable polymers poly (d,l-lactic acid) and poly(e-caprolactone) were prepared by Broz et al. [19]. Figure 7.7 shows the strain-at-failure data across the entire mass fraction range of PLA–PCL blend. This value decreased monotonically as the mass fraction of PLA increases. A rather precipitous drop was seen from 0 to 0.6 followed by flat behavior thereafter. This indicates that even small amounts of glassy PLA are capable of embrittling the PCL matrix; at PLA mass

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Table 7.1 Mechanical properties of 50/50 NR/PMMA blends with graft copolymer [18]. Standard deviation values are given in parentheses. Solution-Cast System Graft Copolymer (%) 0 5 10 15

Melt-Mixed System

Tear Strength (N/mm)

Elongation at Break (%)

Tensile Strength (N/mm2 )

Tear Strength (N/mm)

Elongation at Break (%)

Tensile Strength (N/mm2 )

9.80 (1.62) 20.21 (1.64) 22.40 (1.35) 23.47 (1.26)

322 (33.6) 185 (13.6) 80 (3.28) 132 (4.83)

3.75 (1.06) 9.61 (0.74) 12.37 (0.63) 13.37 (1.16)

9.46 (0.58) 14.75 (0.51) 17.76 (0.32) −

49 (3.52) 31 (3.16) 21 (2) 24 (1.4)

1.25 (0.1) 2.01 (0.12) 3.10 (0.34) 3.14 (0.19)

fraction 0.2 the strain-at-failure has decreased 50% relative to pure PCL. This may be due to blending of the glassy PLA into the PCL matrix or simply to the formation of PLA inclusions in the blend with some interfacial adhesion. Flat part of the curve at PLA mass fractions at and above 0.6 was consistent with the formation of a continuous PLA matrix; PLA is glassy and the strain-at-failure is expected to be insensitive to strain because it depends on molecular parameters such as the free volume in the matrix phase. If this hypothesis were true, it suggests that there is very little mixing between PLA and PCL in this composition range; otherwise, the PCL would have been expected to have a toughening effect on the blend. Therefore, it appears that there may be some association of PLA and PCL at low PLA mass fraction but at higher PLA contents there is little or no reinforcement due to blending.

0.40 0.35

Strain-at-failure

0.30 0.25 0.20 0.15 0.10 0.05 0.00 0.0

0.2

0.4

0.6

0.8

1.0

Mass fraction PLA

Figure 7.7 Plot of strain-at-failure across the composition range for the PLA/PCL blends. Reprinted from [19]. Copyright (2003) with permission from Elsevier.

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40

Yield Stress (MPa)

35 30 25 20 15 10 0.0

0.2

0.4

0.6

0.8

1.0

Mass fraction PLA

Figure 7.8 Plot of yield stress as a function of composition in PLA–PCL blend. Reprinted from [19]. Copyright (2003) with permission from Elsevier.

In Figure 7.8, the yield stress data as a function of composition are shown. The yield stress is insensitive to composition from PLA mass fraction 0–0.4 then increased linearly up to 1.0. This linear dependence of yield stress at compositions ranging from pure PLA down to PLA mass fraction 0.6 suggests PCL blending in this regime simply dilutes the PLA matrix like the presence of voids in the material, reducing the total stress necessary to fracture the samples. The flat trend at compositions ranging from pure PCL to PLA mass fraction 0.4 suggests there must be some reinforcement due to interactions between PLA and PCL, otherwise a similar reduction in yield stress in blended samples would be observed as in the PLA-rich samples. Above a threshold PLA mass fraction of 0.4, the modulus and ultimate tensile strength increased almost linearly as a function of composition. This threshold may be due to strengthening of the blend interface in this regime. George et al. studied the effect of the blend composition on mechanical properties of SBR/NR blends [20]. Mechanical properties of homopolymers and blends are given in Table 7.2. The SBR/NR blends are denoted by N0 , N30 , N50 , N70 , and N100 , where the subscripts denote the weight percent of NR in them. The properties such as tensile strength and elongation at break were increased from N0 to N100 . Mechanical strength of SBR increased upon blending it with NR. This is definitely associated with the strain-induced crystallization of NR. The Young’s modulus decreased from N0 to N100 , which indicates that the initial stretching of SBR and the blend with the higher SBR content requires higher stress. In the swollen state, there is overall reduction in the magnitude of all the mechanical properties. The tensile behavior of the swollen specimens is governed by two types of relaxation mechanisms: the intramolecular motions of segments and the molecular motions involving the adjustments and shifting of chain entanglements. In the equilibrium swollen state, the rubber–solvent interaction is maximum and the rubber–rubber interaction is minimum. This gives rise to the abrupt decrease of the tensile properties of the swollen samples. The mechanical properties of the deswollen samples showed an improvement compared to the unswollen samples. This might be due to increase in the interchain interaction after a sorption–desorption process. The mechanical properties of the blends were also intermediate to those of the homopolymers. The properties of the swollen samples were largely reduced due to the high rubber–solvent interaction during swelling.

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Table 7.2 Mechanical properties of SBR/NR blends: N0 (SBR100); N30 (SBR70/NR30); N50 (SBR50/NR50); N70 (SBR30/NR70); N100 (NR100) [20].

Sample

System

Unswollen

N0 N30 N50 N70 N100 N0 N30 N50 N70 N100 N0 N30 N50 N70 N100

Swollen

Deswollen

7.2.1.4

Secant Modulus (MPa)

Tensile Strength (MPa)

Elongation at Break (%)

Young’s Modulus (MPa)

M100

M200

M100

1.82 4.59 5.24 6.09 6.63 0.70 0.90 1.04 0.64 0.54 1.89 4.62 5.40 6.25 6.81

391 862 869 1007 1069 225 291 342 362 390 398 870 892 1012 1080

1.36 1.22 0.80 0.62 0.35 0.68 0.49 0.59 0.43 0.32 1.72 1.545 1.69 0.603 0.976

0.43 0.62 0.66 0.64 0.37 0.43 0.38 0.36 0.27 0.20 0.883 0.749 0.814 0.959 0.625

1.05 1.01 0.98 0.96 0.58

1.24 1.35 1.34 1.27 0.79

0.66 0.63 0.42 0.32 1.119 1.159 1.343 1.04

0.90 0.42 1.524 1.60 1.40

Influence of crosslinking on mechanical properties

The degree of crosslinking has a strong influence on the mechanical properties of rubber/rubber, rubber/thermoplastic and other blends where one or both phases could be crosslinked. Semba et al. added dicumyl peroxide (DCP) to PLA/PCL binary blend to induce chemical crosslinking, to generate a high performance material [21]. The effects of crosslinking on the mechanical properties of polylactic acid/polycaprolactone blends were studied in detail. There were clear differences at the fracture surfaces obtained after tensile testing, as shown in Figure 7.9. Many dropout traces were observed in the sample without DCP, which decreased with increasing DCP content. The dropout traces of the samples with 0.2 and 0.3 phr of DCP were very few, which is an indication of the good interfacial adhesion between PLA and PCL phases. These results were in agreement with the observed ultimate strain. Variation of tensile properties with DCP content is shown in Figure 7.10. 7.2.1.5

Influence of Compatibilization on Mechanical Properties

Compatibilization is defined as the process of modification of the interfacial properties of the immiscible blend, leading to the creation of a polymer alloy which is an immiscible polymer blend, having a modified interface and/or morphology. When two immiscible polymers are blended without compatibilization, one generally obtains a mixture with a coarse, unstable morphology coupled with poor interfacial adhesion between the phases. As a result, the blends exhibit inferior physical properties to those of individual polymers. A poor interfacial adhesion results in an immature stress transfer which cannot prevent cracks initiation at the interface leading to catastrophic failure. Both theories and experiments support the role of compatibilizer in multiphase polymer systems. It has been assumed that compatibilizers suppress the droplet coalescence in immiscible polymer blends by preventing droplets from approaching each other [22–24]. Compatibilization of blends is broadly classified into two types: Physical compatibilization and reactive compatibilization. Physical

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Figure 7.9 Tensile fracture surface of 70/30 blends. (a) DCP0; (b) DCP0.1; (c) DCP0.2; (d) DCP0.3. Reprinted from [21]. Copyright (2006) with permission from John Wiley & Sons.

compatibilization includes the addition of pre-synthesized copolymer, block (di-,tri-or tapered) copolymer, graft copolymer, random copolymer, gradient copolymer into the immiscible blend (Figure 7.11). In reactive compatibilization the copolymer is formed in situ by a chemical reaction during the extrusion process during the establishment of the immiscible phase morphology. Reactive compatibilization involves a heterogeneous reaction across a phase boundary. Such a reaction is limited by the interfacial volume available at this phase boundary. For incompatible blends containing at least one semi-crystalline component, the final tensile properties are determined by two competing factors: the increase in compatibility due to the presence of more crystalline component and the extent of compatibility between the component polymers. The former is the property determining factor at low strain level and the latter determines the properties at high strain level. The effects of compatibilization on the mechanical behavior of the PP/HDPE blends were studied [26]. Stress-strain behavior of uncompatibilized PP/HDPE blends is demonstrated in Figure 7.12. These curves will also give an approximate indication of the maximum tensile strength (σ m ), elongation at break (Eb ) and Young’s modulus (E). The σ m and Eb showed negative deviation from the additivity line indicating that the blends are highly incompatible. Interestingly, it is found that Young’s modulus experienced a synergism. Compatibility of the

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Figure 7.10 Tensile modulus, strength, and ultimate strain of injection moldings. (a) tensile modulus; (b) tensile strength; (c) ultimate strain; (•) neat PLA; () E70/30; (×) neat PCL. Reprinted from [21]. Copyright (2006) with permission from John Wiley & Sons.

blends does not play a major role in the case of Young’s modulus since it is measured at low strain level. The synergism can be explained in terms of interfacial deformation of the blends. It is believed that during crystallization of the matrix, deformation of the dispersed particle occurs and this results in the deformation of the interface in polymer blends. Thus during the crystallization of the blend where PP is the matrix phase, solidification of PP occurs in the presence of HDPE melt which constitutes the dispersed phase. During this process, molten HDPE flows into the region between PP sperulates growing near the interface. This results in the deformation of the interface between PP and HDPE, and the deformation ends with completion of crystallization sperulates, when all the PP melt is converted to sperulates. The net result is an increase in interfacial area. However, the formed interfaces are so weak that they can transfer stresses only at very low strain levels. This is the reason for the synergism in Young’s modulus, which is measured at low strain levels.

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Figure 7.11 Schematic connecting chain at an interface. (a) diblock copolymers; (b) end grafted chains; (c) triblock copolymers; (d) multiply grafted chains; (e) random copolymer. Reprinted from [25]. Copyright (2002) with permission from Springer.

The effect of compatibilization on the σ m of 80/20 blends is given in Figure 7.13. Ethylene propylene diene terpolymers (EPDM) with three different ethylene /propylene(E/P) ratio were used as the compatibilizers. The σ m of the blends increased by the incorporation of compatibilizers irrespective of the difference in their symmetry in terms of monomer fraction. However, it is also seen that the σ m increased with increase in the amount of compatibilizer, reached maximum at 5 wt% of the compatibilizer and beyond this almost leveling off in strength is observed. This is in good agreement with the morphology of the blends, which revealed that 5wt% compatibilizer concentration was the interfacial saturation point (CMC) beyond which no effective compatibilization took place.

35 30

Stress (MPa)

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H0 H20 H50 H80 H100

10 5 0 0

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20

Strain (%)

Figure 7.12

Stress-strain behavior of uncompatibilized PP/HDPE blends [26].

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Ultimate tensile strength (MPa)

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32

30

28

26 0

5

10

15

20

Weight % of compatibiliser

Figure 7.13 Effect of compatibilization on the tensile strength of HDPE/PP blends (20 parts HDPE).0.5E, 0.6E, 0.7E indicate 0.5, 0.6 and 0.7 of ethylene content in Ethylene propylene diene terpolymers respectively [26].

On the other hand, elongation at break (Eb) of the blends increased with compatibilizer concentration. Even though there is no true leveling off in elongation, the rate of increase decreased beyond 5w% of the compatibilizer concentration. Further, unlike tensile properties, there is difference in performance of different compatibilizers. The Young’s modulus presented in Figure 7.14, on the other hand, shows a different behavior. As the compatibilizer concentration increased, the Young’s modulus experienced marginal decrease. This is basically due to the combined effect of two phenomenon. First, since Young’s modulus is measured at low strain level, as mentioned earlier, compatibility has no significant role. The presence of compatibilizer indeed increased

Figure 7.14 [26].

Effect of compatibilization on the Young’s modulus of HDPE/PP blend containing 20 parts HDPE

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the compatibility and therefore enhanced most of the mechanical properties of the blends, but did not have an impact on the Young’s modulus. Thus increase in compatibility does not increase the Young’s modulus. There are two possibilities for the lower values of Young’s modulus in the presence of compatibilizer: (i) the addition of compatibilizer may reduce the crystallinity of the individual polymer and thereby decreases the Young’s modulus. This can be ruled out as we have seen compatibilizers have no significant effect in the crystallization behavior of PP and HDPE; (ii) the compatibilizers may increase the softness of the blends and this decreases the Young’s modulus. Since the compatibilizer is a flexible polymer, this is a more probable reason for decrease in Young’s modulus [27, 28]. Kim et al. prepared the blends of PBT and EVA by reactive compatibilization of PBT and EVA by MAH (maleic anhydride) [29]. The formation of PBT-g-EVA copolymer as an in situ compatibilizer identified by the reaction of hydroxyl groups and/or carboxylic groups at the chain ends of PBT and MAH grafted onto EVA in the presence of DCP (dicumyl peroxide) as an initiator. When EVA is blended with PBT, the in situ compatibilizer, i.e., PBT-g-EVA, has obtained from the reaction of MAH grafted onto EVA and the hydroxyl groups and/or carboxylic groups at the chain ends of PBT. The grafting yield and the gel content during the reactive compatibilization processes were higher at higher DCP content. The mechanical properties of crosslinked polymeric materials are usually improved with the degree of crosslinking. As a result it was noticed that the flexural strength of the PBT-EVA-g-MAH blend is apparently affected by the crosslinked components of EVA-g-MAH, while the tensile strength is not. The tensile strength decreased but flexural strength increased with the increasing gel contents and grafting yield. 7.2.1.6

Influence of Fillers on Mechanical Properties

In recent years, new kinds of polymeric materials are emerging, such as polymer blends reinforced with micro- and nanofillers such as glass beads, talc, organically modified clays, nanosilica and carbon nanotubes (CNT), which are attracting immense attention because of their remarkable properties. These new kinds of high-performance materials combine the advantages of polymer blends and the merits of polymer nanocomposites. For multi-component systems, one can envision to move beyond simple dispersion and to design strategies that afford the opportunity to selectively reinforce only one of the polymer phases, or even design nanofillers that can promote compatibilized structures of the phases [30–36]. Rahmatpour et al. prepared blend/clay nanocomposites of 50/50 (wt%) NR/SBR [37]. Figure 7.15 shows the mechanical properties of 50/50 NR/SBR blend/clay nanocomposites. The hardness (shore A) and 100% tensile stress (i.e., stress at 100% elongation) of the nanocomposites are higher than those of clay-free 50/50 NR/SBR blend vulcanizate, and also increase with increasing amounts of clay, up to 6 phr, after which they remained almost constant. Increase in the hardness and 100% tensile stress of nanocomposites relative to the clay-free vulcanizate can be attributed to the layered structure of clay and extremely high interfacial action between the silicate layers (or stacked layers) and rubber matrix. The tensile strength and tensile strain at break of NR/SBR blend also are improved by introducing the clay into the rubber matrix. In addition, the improvement increased by increasing the amount of clay up to 6 phr and then decreased slightly in the nanocomposite containing 10 phr clay. It should be noted that all of the mechanical properties of nanocomposites improve considerably with an increasing clay content, up to 6 phr, and then remain almost constant. In other words, maximum improvement in the mechanical properties of nanocomposites can be achieved only when all of the layered silicates are separated into single layers. Vo and Giannelis have prepared poly (vinylidene fluoride)/nylon-6 (PVDF/N6) blend clay nanocomposite [38]. In their study, they used nanoclay as an alternative means to control interfacial properties. Nanoclays are an attractive alternative to traditional compatibilizers. This polymer blend/clay system shows improved stiffness, strength, and toughness. Two different type of clay (Cloisite 30B and Cloisite 20A) and two different mixing stratagies (one batch and sequential compounding) were used for the blend nanocomposite

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700

9 8

600

7 6

500

5 4 3

Tensile strength Tensile strain

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Tensile strength at break (%)

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Clay amount (phr)

Figure 7.15 Mechanical properties of clay-free 50/50 NR/SBR blend vulcanizate and 50/50 NR/ SBR blend/NaMMT nanocomposites. Reprinted from [37]. Copyright (2008) with permission from Taylor & Francis.

preparation. They found that toughness is related to the domain size and degree of crystallinity of the domain phase. The domain size is controlled by the PVDF/N6 viscosity ratio and interfacial tension which are ultimately determined from the surface modification of the nanoclay and the dispersion of the nanoclay particles. To better understand the kinetics of nanoparticle diffusion on blend properties, a series of experiments were performed where the 30B clay was first compounded with one polymer and then the resulting composite was compounded with the second polymer. The 30B clay interacts more favorably with N6, so in the blend in which PVDF is blended with 30B first and then with N6 (referred to as (PVDF/30B)/N6), it is expected that 30B will migrate to the interface and/or the matrix. For the inverse sequence blend in which 30B is compounded first with N6 and then with PVDF (referred to as (N6/30B)/PVDF), it is expected that the 30B nanoparticles will more likely reside in the matrix. TEM images support the overall exfoliation of clay in N6 for both blends (Figure 7.16). Comparison of the TEM images shows a striking difference among the blends’ domain sizes. The PVDF domains in the (PVDF/30B)/N6 blend are ∼110 nm; the domains are not as small as the one batch blend (∼60 nm) but not as large as the domains of the blend with no clay (∼150 nm). However, the (N6/30B)/PVDF sequential blend exhibits larger domains (∼240 nm) than the blend with no clay. Because of the larger domain size and lack of crystallization suppression in the sequence blends, the mechanical properties of the sequential blends were poorer than those of the one batch blend (Figure 7.17). The (N6/clay)/PVDF sequence blend with large domains was stiff (E) 2.52 (0.06 GPa) but showed no improvement in toughness or strength compared to the blend with no clay. Conversely, the (PVDF/clay)/N6 blend was stiffer (E) 2.59 (0.04 GPa), stronger, and tougher compared to the blend with no clay but was inferior to the one batch 30B blend.

7.2.1.7

Modeling of Mechanical Properties

It is well known that mechanical properties might be used to assess the miscibility in polymer blends through a comparison of experimental results and predictions based on various models. Indeed, the mechanical properties of polymer blends depend on the intermolecular forces, chain stiffness, and molecular symmetry of the individual polymers used to prepare the blend [39]. Furthermore, according to Willemse et al. [40],

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Figure 7.16 TEM micrographs of PVDF/N6 30:70 blend nanocomposites with 5% 30B (a) one batch blend; sequential blend in which (b) clay was compounded first with PVDF [(PVDF/30B)/N6] and (c) clay was compounded first with N6 [(N6/30B)/PVDF]. Reprinted from [38]. Copyright (2007) with permission from American Chemical Society.

tensile modulus of polymer blends is strongly dependent on the composition and morphology of blends, and theoretically, it lies between an upper limit given by the parallel model: M = M1 ϕ1 + M2 ϕ2

(7.1)

where M is the modulus of the blend, M 1 and M 2 are the moduli of the components 1 and 2 respectively, ϕ 1 and ϕ 2 are the volume fraction of the components 1 and 2 respectively. This equation is applicable for models in which the components are arranged parallel to the applied stress.

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Stress (MPa)

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60 40 20 0 0

20

40

60

80

100

120

% Strain

Figure 7.17 Tensile stress-strain curves for PVDF/N6 blend and 30B modified blends. Reprinted from [38]. Copyright (2007) with permission from American Chemical Society.

The lower bound of the modulus holds models in which the components are arranged in series with the applied stress and the equation is: 1/M = ϕ1 /M1 + ϕ2 /M2

(7.2)

M1 /M = (I + Ai Bi ϕ2 )i(l − Bi ϕ2 ) Bi = (M1 /M2 − l)/(M1 /M2 + Ai )

(7.3) (7.4)

According to Halpin-Tsai equation:

In the Halpin-Tsai equation, subscripts 1 and 2 refer to the continuous and dispersed phase respectively. The constant Ai is defined by the morphology of the system. For elastomer domains dispersed in a continuous hard matrix, Ai = 0.66. Moreover, for a cocontinuous system, the modulus has to agree with the Davies model, where the equation is given by: 1/5

E 1/5 = φ1 .E 1/5 + φ2 .E 2

(7.5)

in which Ei and ϕ i are the elastic modulus and thevolume fraction of phase i, respectively. The Coran–Patel model [41], which represents a phenomenological adjustment between the parallel and series models, is E = φ nH (nφs + 1)(EU − E L ) + E L

(7.6)

where EU is the upper bound and EL is the lower bound and n is an adjustable parameter related to the change in phase morphology as a function of H. Kunori and Geil [42] reported that when a strong adhesive force exists between the blend components the dispersed phase will contribute to the strength of the blend. The equation is: σb = σm (1 − Ad ) + σd Ad

(7.7)

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where Ad represents the area occupied by the dispersed phase in the transverse cross section. Considering two possible fracture paths in a blend, the equation can be modified as follows, depending on whether the fracture is through the interface or through the matrix. When the fracture is through the interface: 2/3

2/3

σb = σm (1 − φd ) + σb φd

(7.8)

When the fracture is through the matrix: σb = σm (1 − φd ) + σd φd

(7.9)

where σ b , σ m , and σ d are the properties of the blend, matrix phase, and dispersed phase respectively, and ϕ d is the volume fraction of the dispersed phase. Another important model for perfect adhesion is the Kerner equation [43]. According to this:

E = Ee

φd E d (7 − 5νm )E m +(8 − 10νm )E d + φm /15(1 − νm )

(7.10)

φd E m /(7 − 5νm )E m +(8 − 10νm )E d + φm /15(1 − νm )

where E, Em , and Ed are the respective properties of the blend, continuous phase, and dispersed phase; ϕ d and ϕ m , the volume fractions of the dispersed and continuous phases; and ν m , the Poisson’s ratio of the continuous phase. Varghese et al. [15] applied the models such as the parallel model, series model and Halpin-Tsai equation to predict the mechanical and viscoelastic behavior of NBR/EVA blends. The applicability of these simple models to mechanical and viscoelastic properties is presented in Figure 7.18. In all cases it is seen that the experimental data are close to the parallel model.

(a)

(b)

200 Storage modulus, E ′ (x107 N/m2)

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150

Experimental Parallel Series Helpin–Tsai equation

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100

5

0 0

20

40

60

80

Volume fraction of NBR

Experimental Parallel Series Helpin–Tsai equation

100

50

0

0

20

40

60

80

100

Volume fraction of NBR

Figure 7.18 (a) Applicability of various models on tensile strength of NBR/EVA blends; (b) Applicability of various models on storage modulus of NBR/EVA blends at 10◦ C [15].

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TENSILE STRENGH (MPa)

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0.3

0.5

0.7

1

VOLUME FRACTION OF NR Experimental Kunori

Figure 7.19 of NR [20].

Parallel Kerner

Series

Comparison of experimental tensile strength with theoretical values as a function of volume fraction

Mechanical behavior of SBR/NR blends was modeled [20] using various composite models such as the parallel, the series, the Kerner and the Kunori models. Figure 7.19 shows the theoretical and experimental curves of the tensile strength values of the SBR/NR blend. In N30 and N50 , the experimental values are close to that of the Kunori model. Therefore, it can be concluded that the fracture propagates through the interface rather than through the matrix. In N70 , the experimental value is close to that of parallel model, Therefore, in N70 , it can be concluded that the applied stress distributes equally in two phases. Veenstra et al. [17] proposed a model to predict the moduli of polymer blends with co-continuous morphologies over the complete composition range. This model is obtained by depicting the co-continuous morphology as three orthogonal bars of the first component embedded in a unit cube where the remaining volume is occupied by the second component, leading to a series model of parallel parts and a parallel model of serial-linked parts. A judicial use of these models results in a perfect description of the moduli of co-continuous blends as a function of composition. To visualize co-continuity, the model consists of three orthogonal bars of polymer 1 embedded in a unit cube where the remaining volume is occupied by component 2. Repeating this unit cube in 3D shows that component 2 has the same framework as component 2, i.e., both the components are interconnected. Relations for a series model of parallel parts (Figure 7.20(a) and Eq. (7.11)) and for a parallel model of serial-linked parts (Figure 7.20(b) and Eq. (7.12)] can be derived [44] as: Ec =

(a 4 + 2a 3 b)E 12 + 2(a 3 b + 3a 2 b2 + ab3 )E 1 E 2 + (2ab3 + b4 )E 22 (a 3 + a 2 b + 2ab2 )E 1 + (2a 2 b + ab2 + b3 )E 2

(7.11)

Ed =

a 2 bE 12 + (a 3 + 2ab + b3 )E 1 E 2 + ab2 E 22 bE 1 + a E 2

(7.12)

where a is related to the volume fraction of component 1 by 3a2 2 2a3 = ϕ 1 and b is related to the volume fraction of component 2 by b =1− a.

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Figure 7.20 Three-dimensional models for the calculation of the moduli of co-continuous polymer blends (a and b). Reprinted from [17]. Copyright (2000) with permission from Elsevier.

They compared the Young’s moduli of co-continuous blends of polyethylene (PE) and polystyrene (PS) to the models as proposed in Eqs. (7.11) and (7.12). Again they used Eq. (7.12) in the region (35–45 vol.%) where the soft phase dominates, Eq. (7.11) in the region (50–80 vol.%) where the stiff phase dominates, and an average between both the equations in the intermediate range where neither phase dominates. The linear interpolations result in a stepwise plot. The results can be seen in Figure 7.21. The agreement between the moduli that are predicted by the models and the experimental data for the co-continuous PE/PS blends is again satisfactory. 7.2.2

Dynamic mechanical properties

Dynamic mechanical thermal analysis (DMTA) is another powerful technique to investigate the performance of polymer blends, as it measures responses of a material to cyclic stress. The investigation of dynamic modulus

Ey (MPa)

5000

1000

100 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 Volume fraction PS

Figure 7.21 Young’s moduli of co-continuous PE/PS blends (•) and the predictions (full line) using Eqs. (7.11) and (7.12). For comparison, the parallel (—), series (– – –) and Davies (. . . . . .) model are shown. Reprinted from [17]. Copyright (2000) with permission from Elsevier.

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and damping behavior over a wide range of temperatures and frequencies has proven to be very useful in studying the structural features of polymer blends and the variation of properties with respect to end use applications [45–52]. These rely on structure, crystallinity, extent of crosslinking, etc., which in turn depends on the phase morphology of blends. Note that the dynamic mechanical properties are sensitive not only to different molecular motions but also to various transitions, relaxation processes, structural heterogeneity and the morphology of multiphase systems. Further, the dynamic mechanical properties of polymers give mirror images of their molecular and morphological features. Another important aspect is that polymers are viscoelastic materials, which have some of the characteristics of both viscous liquids and elastic solids. Elastic materials have a capacity to store mechanical energy with no dissipation of energy; on the other hand, a viscous fluid in a non-hydrostatic stress state has a capacity for dissipating energy but none for storing it. When polymeric materials are deformed, part of the energy is stored as potential energy and part is dissipated as heat. The energy dissipated as heat manifests itself as mechanical damping or internal friction. Therefore the interpretations of these properties at molecular level are of great scientific and practical importance in understanding the mechanical behavior of polymers. Dynamic mechanical thermal analysis helps to measure the glass transition temperature of polymers. In addition one can obtain an idea about the storage (dynamic) modulus, loss modulus, and damping behavior (internal friction). The Tg of polymers is a very important parameter since there are profound changes in the physical properties of the polymers –heat capacity, thermal expansion coefficient and modulus occur at this temperature. It should be noted that Tg of the polymer is accompanied by a sharp decrease in stiffness. The stress relaxation modulus decreases and the creep compliance increases by about three orders of magnitude in glass transition region. The loss moduli and loss compliances exhibit a maximum in the glass transition as does tan δ. Any type of interaction between the polymers will give rise to a shift in this maximum and thus dynamic mechanical properties will provide an idea about the extent of interactions between the component polymers in a polymer blend. The effect of different compatibilizers on the dynamic mechanical properties of PS/polybutadiene blends was studied by Joseph [53]. The author reported that storage modulus, loss modulus and tan δ underwent dramatic change in the presence of compatibilizers. For LLDPE/EVA blends, it has been found that compatibilization increased the storage modulus of the system which is due to the dispersion of EVA domains in the LLDPE matrix providing an increased interfacial interaction [54]. Figure 7.22 demonstrates the effect of blend ratio on the storage modulus (E ) of PP/HDPE blends [26]. PP exhibits the maximum and HDPE the minimum storage moduli in the whole temperature range. The E of all blend exhibits intermediate values. As the weight % of HDPE in the blend increases, storage modulus increases. However, it is interesting to compare the Young’s modulus with storage modulus since both are measured under tension. However, the E values offer apparent conflict with the Young’s moduli and are considerably higher than the Young’s moduli. At the same time there is no synergism in the E values. The effect of blend ratio on the Tg of PP are given in Figure 7.23. Separate peaks suggest that the HDPE/PP blend is highly incompatible at all combinations. The vulcanization of the rubbery phase during mixing has been investigated as a way to improve the physical properties of several thermoplastic elastomers based on rubber/plastic blends. The change in morphology that occurs during dynamic vulcanization is schematically represented in Figure 7.24. During dynamic vulcanization, a co-continuous morphology may be transferred to a matrix and dispersed-phase morphology, there may be some possibility of phase inversion, or the crosslinked rubber phase may become finely and uniformly dispersed in the plastic matrix. During the process of dynamic vulcanization, the viscosity of the rubber phase increases because of crosslinking, and the rubber domains can no longer be sufficiently deformed by the local shear stress and are eventually broken down into small droplets. The variations of G , G and tan δ of 70/30 HDPE/ EVA blends crosslinked with 0.5 and 1.5% DCP (dicumyl peroxide) were studied by John et al. [55]. The addition of peroxide to 70/30 HDPE/EVA blends leads to

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273

The variation of storage modulus of PP, HDPE and their blends as a function of temperature [26].

an increase in G and G (Figure 7.25). This increase is more pronounced at the higher DCP contents. At temperatures greater than 100◦ C, 1.5% DCP leads to a decrease in G . The α- and γ -relaxation temperatures of HDPE are not influenced significantly by the addition of DCP. The relaxation of EVA at about –20◦ C is slightly changed by the addition of 0.5% DCP. The addition of 1.5% DCP leads to significant increase in the relaxation temperature of EVA by 3.6 K, which indicates the predominant crosslinking of the EVA phase. In addition, the intensity of the tan δ peak corresponding to the EVA content increases with the increase in the DCP content. In this case also, the blends show the presence of two peaks corresponding to Tg’s of HDPE and EVA, which indicates the immiscibility of the system. The peak widths at the half-heights of the

Figure 7.23

Effect of blend ratio on the variation of tan δ as a function of temperature in HDPE/PP blend [26].

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Figure 7.24

Schematic representation of the dynamic vulcanization [55].

10

Storage modulus G’(Pa)

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HDPE/EVA 70/30 70/30 + 0.5% DCP 70/30 + 1.5% DCP –150 –100

–50

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Temperature (°C)

Figure 7.25

Dependence of G on the temperature of 70/30 HDPE/EVA dynamically crosslinked blends [55].

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Table 7.3 Storage modulus of LDPE, POE, LDPE/POE blends, and LDPE/POE/clay nanocomposites [56]. Storage modulus/GPa Sample No. 1 15 5 8 14 13 11 12 9 10 2 16

−140◦ C

−100◦ C

−50◦ C

0◦ C

25◦ C

3.57 3.35 3.51 3.73 3.44 4.29 3.37 3.74 2.51 2.99 3.49 4.08

2.40 2.38 2.35 2.49 2.40 2.90 2.23 2.50 1.66 2.02 2.12 2.78

1.67 1.72 1.55 1.68 1.61 2.02 1.32 1.48 0.85 1.08 0.84 1.18

0.65 0.65 0.36 0.39 0.36 0.54 0.17 0.20 0.06 0.08 0.03 0.04

0.34 0.33 0.19 0.20 0.19 0.28 0.09 0.10 0.03 0.04 0.02 0.02

70/30 HDPE/EVA blend and the 70/30 HDPE/EVA blends crosslinked with 0.5 and 1.5% DCP were 31, 33, and 37◦ C, respectively. This increase in the peak width or broadening of peaks upon crosslinking indicates that dynamic crosslinking promotes interfacial bonding between the phases, which may arise because of the crosslinking of HDPE and EVA phases. Dynamic mechanical properties of low-density polyethylene (LDPE)/ethylene–octene copolymer (POE)/organomontmorillonite (OMMT) nanocomposites were studied by Baghaei et al. [56]. The storage modulus of the LDPE/POE blends and nanocomposites at different temperature regions are shown in Table 7.3. As can be seen, the storage modulus of the nanocomposites are higher than their neat blends, indicating the reinforcing role of clay. OMMT anchors at different positions in the matrix, thus restricting the movement of the chains. As seen from Table 7.3, at temperatures higher than glass transition temperature (Tg) of POE (–50 ◦ C) storage modulus of the POE/OMMT (95/5) nanocomposite is about 50% higher than that of the neat POE. The enhancement of storage modulus strongly depends on the aspect ratio of the dispersed clay layers and the intercalation of the polymer chains inside the clay matrix. When the polymer matrix is reinforced with a rigid filler the polymer interface adjacent to the clay particle is highly restrained mechanically. Active surface area of the filler increases due to the intercalation of the polymer chains inside the clay galleries. Polymer chains inside the clay galleries are immobilized and the effective immobilization of these chains is responsible for the enhancement of the hydrodynamic storage modulus 7.2.2.1

Modeling of Dynamic Mechanical Properties

Kim et al. applied parallel-series model, Davis model and Coran–Patel model to methyl acrylate (MA) and 2-acrylamido-2-methyl propane sulfonic acid (AMPS abbreviated as AP) modified polyacrylonitrile/cellulose acetate (MA–PAN–CA and AP-PAN-CA) blends [57]. MA–PAN–CA blends (Figure 7.26) shows negative deviation from the Davies model, whereas AP–PAN–CA blends (Figure 7.27) are relatively well fitted with the model. These results imply that interfacial interactions are poor in MA–PAN–CA blends and relatively good in AP–PAN–CA blends, and the results are consistent with dynamic mechanical data and morphology.

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9.9 Log E(dyne/cm2)

9.6

9.3

9.0

20

0

40 60 CA content (wt%)

80

100

Figure 7.26 Complex moduli of MA–PAN–CA blends at 80◦ C, and comparison with upper–lower bound model (solid line), Davis model (bold solid line), and Coran–Patel model (dashed line). Reprinted from [56]. Copyright (2009) with permission from Springer.

For MA–PAN– CA blends, the 80/20 blend is fitted with n = 3 of the Coran–Patel model (Eq 7.6), 40/60 and 60/40 blends with n = 3 – 4, and the 20/80 blend with n = 5, and the experimental values are generally well fitted with n = 3 – 4. However, in the AP–PAN–CA blend, experimental values are well fitted with n = 1 – 2. (n − 1)/n indicates the center of the H range where phase inversion or transition occurs. Thus, it can be noted that phase inversion takes place at H = 0.67 – 0.75 (MA–PAN–CA) and 0.5 – 0.67 (AP–PAN–CA blend). These results agree well with morphology studies. 7.2.3

Impact Properties

Extensive research and development work have been carried out to formulate polymers with high impact resistance. Both rubbers and engineering thermoplastics are incorporated into brittle plastics to improve their 10.2

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CA content (wt%)

Figure 7.27 Complex moduli of AP–PAN–CA blends at 80◦ C, and comparison with upper–lower bound model (solid line), Davis model (bold solid line), and Coran–Patel model (dashed line). Reprinted from [56]. Copyright (2009) with permission from Springer.

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impact properties [58–60]. Impact modification by rubber toughening involves the incorporation of small amounts of rubber (mostly between 3 to 20 vol%) in rigid polymer materials such as glassy thermoplastics, semicrystalline thermoplastic, and thermosets in order to enhance their fracture resistance. The impact resistance of rubber-toughened polymers is strongly dependent upon the concentration, size and size distribution of the rubber particles [61]. The role of particle size distribution has been studied by Wu [62]. According to Wu, a wide particle size distribution is disadvantageous for the impact behavior of rubber-modified blends with pseudoductile matrices because the interparticle distance increases with the increasing size of polydispersity at a give rubber fraction. According to Bucknall and Paul, in high-performance blends, toughness reaches a maximum at particle diameters of 0.2–0.4 μm [63]. The critical interparticle distance (IPD) is related to the size of the rubber particles and the rubber volume fraction, ϕ r , by the following relation. IPD = d[k(/6ϕr )1/3 l]

(7.13)

where the parameter k is a measure of the packing arrangement of a particular lattice and d is the average diameter of the dispersed rubber particle. A bimodal distribution can lead to an enhancement of the toughness. With respect to the role of interfacial adhesion, a distinction has been made between matrices deforming by multiple shear yielding and those deforming by multiple crazing. For matrices deformed by multiple crazes, good adhesion between the rubber particle and matrix is required because the rubber particles must act as effective craze stoppers. For matrices deformed by shear yielding as a result of internal rubber cavitation, a good interfacial strength is not required. The particle size required for the cavitation, d, is given by the following equation: D = 12(γr + sc )/K r 4/3

(7.14)

where is the volume strain, Kr is the rubber bulk modulus, sc is the surface energy per unit area, and γ r is contribution from van der Waals surface tension. Impact properties measurements of PA 6/EPR showed that notched impact strength curves are strongly influenced by the particle size, as shown in Figure 7.28 [8]. Borggreve and Gaymans studied the effect of the coupling agent on the impact behavior of nylon 6EPDM (ethylene propylene diene monomer) rubber. The rubber was grafted with various amounts of maleic anhydride (MA) with the aid of a peroxide. The MA grafted onto the rubber was found to react with the nylon during the blending process. With the MA-grafted rubbers, a much finer dispersion could be obtained. However, the concentration of the coupling agent, within the range 0.13 to 0.89 wt% grafted onto the rubber, has hardly any influence on either the dispersion process or the impact behavior of the blends [64]. The influence of mixing conditions on mechanical properties of low density polyethylene–polystyrene blends were studied by Vasilenko et al. [65]. LDPE–PS blends were obtained by two methods: (a) by mechanical mixing in a melt at 160◦ C in a closed mixer, and (b) by the (high temperature shear deformation) HTSD method in a rotary disperser. For PS, the impact strength is one of the most important mechanical characteristics. It can be improved by introducing LDPE. Figure 7.29 shows that the impact strength of the LDPE–PS blends prepared by the HTSD method exceeds by a factor of 2.0–2.5 the impact strength of the samples obtained by mixing in the melt. Semba et al. added dicumyl peroxide (DCP) to PLA/PCL binary blend to induce chemical crosslinking, to generate a high performance material [21]. The effect of crosslinking on impact properties of polylactic acid/polycaprolactone blends were studied in detail. The impact strength of PLA maintained a constant value at all DCP contents in both tests. On the other hand, the impact strength of PLA/PCL (70/30) showed a higher

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Figure 7.28 [8].

Noted Izod measurements as a function of the weight average particle size for PA6/EPR (20 vol%)

value, which increased with increasing DCP content at low DCP concentrations. In the case of the normal Izod impact test, the PLA/PCL containing 0.3 phr of DCP showed a value that was 2.5 times the value of the PLA. In the case of the reverse Izod impact test, the impact strength of both samples showed similar values from 0 to 0.2 phr of DCP contents. On the other hand, PLA/PCL blend containing 0.3 phr showed a higher value that was 2.5 times those of the other materials. Based on those observations, it is obvious that DCP addition significantly alters the impact properties of this blend system.

50 2

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LDPE content, wt %

Figure 7.29 Relative impact strength of the LDPE–PS blend vs. the content of LDPE with respect to the impact strength of PS: (1) HTSD mixing and (2) mixing in the melt. Reprinted from [65]. Copyright (2009) with permission from Springer.

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IZOD IMPACT STRENGTH (J/m)

1200 1000 20% SEBS-g-MA 800 600 400 20% EPR-g-MA 200 0

0% Rubber 0 NYLON 6

20

40

60

% WEIGHT

80

100 PP

Figure 7.30 Izod impact strength of nylon 6/polypropylcnc blends modified with 20% EPR-g-MA or SEBS-g-MA; the composition shown on the abscissa is the percentage of polypropylene relative to nylon 6 in the blend on a rubber-free basis. The dashed lines correspond to the properties of unmodified binary blends. Reprinted from [66]. Copyright (1995) with permission from Elsevier.

Paul and co-workers studied the effect of the two maleated rubbers EPR (ethylene-propylene random copolymer)-g-MA and SEBS (styrene-ethylene/butylenestyrene triblock copolymer)-g-MA, containing equivalent amounts of maleic anhydride, on the mechanical properties of nylon 6/PP blends [66]. Figure 7.30 shows that the Izod impact strength is greatly improved by addition of 20% by weight of the maleated rubbers, equivalent to 25.1% and 23.8% by volume of EPR-g-MA and SEBS-g-MA, respectively. The improvement in impact strength is relatively independent of the type of rubber used. High impact strengths were obtained in the range 0–50% of polypropylene. Beyond 50% polypropylene, the impact strength is lower: this reduction appears to occur at the point where the polypropylene becomes the continuous phase. The impact strength shows a maximum at about 20–25% polypropylene for both rubbers. Toughening of thermoset for improvement of crack resistance has been the subject of intense research interest during the last two decades. Reactive liquid rubbers and functionally terminated engineering thermoplastics are the widely used toughening agents for the epoxy thermoset. Carboxyl-terminated butadiene acrylonitrile (CTBN) copolymer has been used to modify the aromatic amine cured epoxy resin and anhydride cured epoxy [67, 68]. Appreciable improvements in impact strength were observed in the prepared blend systems (Figure 7.31). Impact behavior of the cured epoxy could be explained based on the two-phase nature of the system. According to Bucknall [61] the rubber particles were considered to bridge the crack as it propagates through the material. Thus, the rubber particles were able to prevent the crack growing to a catastrophic size. The increase in toughness was due to the amount of elastic energy stored in the rubber particles during stretching. Thus, the deformation of the rubber particles in the matrix seemed to be responsible for the enhanced stress transfer and hence impact resistance. Shear yielding of the matrix was another reasonable mechanism that might be operating. According to Newman and Strella [69] the principle function of the rubber particle was to produce sufficient triaxial tension in the matrix so as to increase the local free volume and hence enable extensive shear yielding of the matrix. Thus, crack building of rubber particles along with shear yielding was the main toughening mechanism and enhancement of impact behavior. 7.2.4

Nanostructured Polymer Blends

The idealized morphology of nanostructured polymer blend systems is characterized by the molecular level dispersion of the phases that leads to a considerable enhancement in the mechanical properties, especially

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Toughness, J

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CTBN, %

Figure 7.31 Variation of impact strength and toughness with CTBN content in epoxy/CTBN blend cured with aromatic amine. Reprinted from [67]. Copyright (2007) with permission from Elsevier.

the modulus. In recent years, several strategies have been developed from well-defined and predictable multicomponent polymer structures with phase separation at nanoscale [70]. The most straightforward approach is to use linear block copolymers with two components, A and B. Significantly different morphologies and mechanical properties were generated when replacing a mixture of two almost immiscible linear polymers (PS and PMMA) by the corresponding block copolymer (PS-b-PMMA), to modify an epoxy resin. As shown in Figure 7.32, the morphology obtained for the BC/E (block copolymer-epoxy) blend was quite different than the PS/PMMA/E blends. The morphology of BC/E blend is characterized by the presence of bicontinuous phases, a fact which is highly desirable when the aim is to increase the fracture toughness of the thermoplastic/thermoset blend. Phase separation of the more miscible block induced by polymerization led to the generation of a bicontinuous thermoplastic/thermoset structure exhibiting the desired decrease in yield stress which is necessary

Figure 7.32 SEM micrographs of (a) PS/PMMA/E blends; (b) BC/E blends. Reprinted from [71]. Copyright (2003) with permission from American Chemical Society.

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Table 7.4 Elastic modulus (E) and yield stress (σ Y ) of the neat epoxy matrix and different blends [60]. Blend Neat E PS/E PMMA/E PS/PMMA/E BC/E

E(GPa)

δγ (MPa)

2.65±0.06 2.54±0.01 2.70±0.05 2.64±0.01 2.14±0.02

103.3±0.4 97.1±1.2 98.0±0.5 97.2±1.0 90.0±1.8

for toughening purposes. Values of the elastic modulus (E) and yield stress (σ Y ) of the neat epoxy matrix and different blends are shown in Table 7.4. Ritzenthaler et al. studied the mechanical properties of nanostructured blends of polystyrene-blockpolybutadiene-block-poly-(methyl methacrylate) (SBM) copolymer triblock and epoxy as a function of composition and concentration of SBM [71]. The addition of SBM triblocks can then be expected to be a powerful way of improving the poor fracture resistance of epoxy thermosets. Two SBM copolymers were used: S27 22 B9 M69- SB21 and S14 12 B18 M70- SB10 (S27 22 B9 M69- SB21 represent a copolymer as received and composed of S27 22 B9 M69 pure triblock copolymer and 21 wt % of SB diblock. The numbers 22, 9, and 69 represent the weight of the respective PS, PB, and PMMA blocks; 27 is the molar mass of the PS block in kg mol-1). Figure 7.34 shows the KIc values obtained for different concentrations of SBM up to 50 wt %. It is worth noting that the continuous matrix always remains the epoxy-diamine component whatever the SBM concentration in the investigated range. This constitutes a major difference and advantage compared to the classical use of low molar mass reactive rubbers, for which phase inversion occurs when the concentration is higher than 20 wt %. For both SBM copolymers, the toughness very significantly increases with the triblocks (a)

(b)

Figure 7.33 Transmission electron micrographs DGEBA-MCDEA epoxy systems containing 30 wt %. (a) S27 22 B9 M69- SB21; (b)S14 12 B18 M70- SB10. Reprinted from [71]. Copyright (2003) with permission from American Chemical Society.

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KIC (MPa.m1/2)

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1

0.5

0 0

10

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30 40 SBM wt %

50

60

70

Figure 7.34 Influence of as-received SBM concentrations and compositions on material toughness: (*) and neat DGEBA-MCDEA, (•) S27 22 B9 M69- SB21, () S14 12 B18 M70- SB10 as toughner agents at different concentrations. Reprinted from [71]. Copyright (2003) with permission from American Chemical Society.

concentration. Until 10 wt % of SBM, the toughness is slightly higher for the blends based on the SBM containing the longer PB block. Alternatively, for higher concentrations, blends based on S27 22 B9 M69- SB21 exhibit higher toughness, suggesting that the ‘raspberry’ morphology (Figure 7.33a) is more efficient in improving mechanical properties than the ‘onion’(Figure 7.33b) morphology. Recently, Thio et al. have reported the behavior of poly-(ethylene oxide)-b-poly (hexylene oxide) (PEO–PHO) diblocks in phenol novolac cured BADGE [72]. The highly non-polar PHO block is immiscible with BADGE even at low molecular weights. Wormlike micelles were generated by mixing a vesicle forming diblock (9 wt% PEO) and a spherical micelle forming diblock (44 wt% PEO).Here the wormlike morphology, shown in Figure 7.35, was found to give the best improvement in K1c (∼6×), but vesicles were found to give greater improvements than spherical micelles (∼3.5× and ∼1.75×, respectively).

7.3

Interpenetrating Polymer Networks (IPNs)

According to Sperling, the IPNs are defined as a combination of two or more polymers in network form that are synthesized in juxtaposition [73]. For thermoset IPNs, the related networks are chemically crosslinked. The thickness of the network strands varies from 10 to 1000 nm [74]. The IPN structure offers unique possibilities for composites with respect to fiber to matrix adhesion as schematically depicted in Figure 7.36. Thermoset IPNs are prepared sequentially or simultaneously. In the sequential case, the first crosslinked polymer network is swollen by the monomer of the second polymer that is polymerized and/or crosslinked afterward. The simultaneous IPNS are prepared by crosslinking of two or more monomer systems at the same time. The related reactions should be combination of free-radical induced polymerization and ionic homo-polymerization reactions. The cross-reactions between the crosslinked networks result in grafted IPNs. Thermoset IPNS have been mostly used as tough, impact modified polymers, matrices in composites and as sound and vibration damping materials [73, 76, 77]. Several reports are now available on thermoset

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Figure 7.35 TEM image showing wormlike micelles of PEO–PHO in a BADGE+ PN system. Stained with RuO4. Reprinted from [72]. Copyright (2006) with permission from American Chemical Society.

IPNs, usually chain and step polymerization crosslinking methods are combined. As mentioned before, chain polymerization covers free-radical induced crosslinking of polymers and monomers bearing two or more double bonds along their chains. This is the case with unsaturated polyester (UPs) with VE resins, and crosslinking of vinyl compounds possessing one double bond with two or multifunctional monomers of similar structure. Table 7.5 lists the mechanical and fracture mechanical properties of VE/EP hybrids, the EP phase of which was amine cured. As expected, all IPNS are very sensitive to post curing, which results in a substantial increase in the Tg. This is associated with decrease in fracture energy (Gc ). It was found that VE/VP hybrids, in which the idealized network of the EP phase is composed of aliphatic and cycloaliphatic units, show outstanding toughness values [78].

Figure 7.36 Scheme of the intermittent fiber bonding achieved via matrix organization (i.e., IPN structure). Reprinted from [75]. Copyright (2003) with permission from Springer.

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Table 7.5 Tensile mechanical properties, fracture energy (Gc), and glass transition temperature (Tg) of VE/EP (1:1) systems the EP component of which was cured by polyaddition reaction with diamines [78].

Composition VE/AI-EP+Cal-Am VE/Cal-EP+Al-Am VE

Maximum curing temperature (◦ C)

Tensile strength (MPa)

Ultimate tensile Strain (%)

Young’s modulus (MPa)

Gc (kJ/m2)

Tg (◦ C)

150 200 150 200 150 200

50 60 14 46 48 52

5.5 4.5 38.5 5.1 2.8 2.1

2600 2700 700 2900 3200 3400

3.7 1.1 7.3 5.2 0.54 0.45

81 128 57 87 157 160

Kim et al. evaluated tensile and shear moduli of polyurethane-poly (methy1 methacrylate) interpenetrating polymer networks (IPNs) [79]. The strain rate dependence of tensile modulus is shown in Figure 7.37. The modulus-strain rate curves appear to lie in three groups. The low modulus group includes elastomeric samples where the polyurethane phase is continuous, the medium modulus group includes leathery samples where the phases are in the inversion process (each phase is locally connected), and the high modulus group includes glassy samples where the poly(methy1 methacrylate) phase is continuous. The 75/25% polyurethanepoly (methy1 methacrylate) IPN shows little modulus reinforcement over the pure polyurethane samples. Shear storage modulus, G , vs. temperature plot (Figure 7.38) for polyurethane-poly (methy1methacrylate) IPNs also reveals three general group. The 75/25% and 85/15% polyurethane-poly(methy1 methacrylate) IPN’s show a slightly reduced modulus compared with the pure polyurethane sample at room temperature.

Young’s (dyne/cm2)

1010

UC100 UC75MC25 UC60MC40 UC50MC50 UC40MC60 UC25MC75 UC15MC85 MC100

109

108 0.001

0.1 0.01 Strain Rate (min−1)

1.0

Figure 7.37 Young’s modulus vs. strain rate for the polyurethanepoly(methy1 methacrylate) IPN’s at 23◦ C. Reprinted from [79]. Copyright (1977) with permission from American Chemical Society.

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1010 109 108 107 106

−90

UC100 UC85MC15 UC75MC25 UC60MC40 UC50MC50

−50

−10

UC40MC60 UC25MC75 UC15MC85 MC100

30

70

110

Temperature (°C)

Figure 7.38 Shear storage modulus, G’, vs. temperature for the polyurethane-poly(methy1 methacrylate) IPNs. Reprinted from [79]. Copyright (1977) with permission from American Chemical Society.

Modulus studies revealed that IPN has the lowest modulus at room temperature. This is probably due to crystallization of the linear polyurethane. The above comparison seems to indicate that the reduction of modulus is related to the interaction and/or interpenetration of the phases. One possible explanation would be that the partial crystallinity in the polyurethane phase (shown earlier to be present, even in the crosslinked polymer) was reduced due to the interpenetration of poly (methy1methacrylate) chains into the polyurethane phase. In the pseudo-IPNs and linear polyblend, the crystallinity of the polyurethane phase was not reduced since there was no interpenetration (interlocking of chains), thus yielding an increased modulus over the pure polyurethane. The increase in modulus in the temperature range of –10 to 30◦ C is due, most likely, to crystallization of the linear polyurethane caused by the slow heating. Moreover, the shift in Tg to higher temperature (most prominent in the full IPNs) also indicates interpenetration. In another study by Huelck et al., the physical and mechanical properties of poly (ethy1 acrylate)-poly (styrene-co-methyl methacrylate), PEAP(S-co-MMA), interpenetrating polymer networks (IPN’s) have been investigated [80]. Stress–strain and tensile data show that the work to break as well as the actual tensile values of the samples steadily increase as the amount of plastic component is increased in the elastomer-rich materials. Mechanical properties of the IPNs based on PI/PMMA were correlated to the morphology of the system [81]. Tensile strength of the IPN samples are plotted as a function of the overall PMMA content as shown in Figure 7.39. The IPNs showed a gradual change from a rubbery to a plastic nature with the increase in concentration of PMMA. Tensile strength and modulus were strongly dependent upon the morphology and showed a dramatic change near the 50/50 composition where the system developed a nanostructured morphology (Figure 7.40). In particular, the tensile strength and modulus showed a lower value in the sea-island region, and a distinct increase when the morphology changed to compact, ordered and smaller PMMA domains having size less than 90 nm and an immediate decrease on the appearance of dual phase morphology. When the PMMA concentration was higher than 30 wt%, tensile strength increased significantly. Elongation at break was found to decrease with the increasing PMMA content due to the low elongation and brittleness of the PMMA phase. These results suggest that the nanostructured morphology obtained around 50/50 composition was the key factor that contributed to the mechanical performance of these IPNs.

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20 Tensile strength MPa

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10

5

0 0

10

20

30

40

50

60

Weight percentage of PMMA

Figure 7.39

Effect of wt % of PMMA on tensile strength of full IPNs [81].

The introduction of crosslinks in PMMA phase was found to improve the mechanical performance of the material. This was due to the higher degree of interpenetration between the PI and PMMA chains during IPN formation. Thus, some degree of enforced miscibility was obtained, which enabled the material to have good mechanical properties. Higher amount of PMMA (more than 55%) was found to reduce the tensile strength. Effect of crosslink density of PI on the mechanical properties was also studied by preparing IPNs with highly crosslinked PI samples. A considerable decrease in tensile strength was observed with the increase in crosslinking degree of the PI phase. However, the effect of crosslink density of PMMA on the mechanical performance of full IPNs was different from that of the PI phase. The lower crosslink in PMMA favors the

Figure 7.40

TEM image of IPN s with nanometre-sized morphology [81].

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413K 383K 353K 333K 303K s f

3260

3280

s

293K

f

193K

3300 3320 Field (G)

3340

Figure 7.41 ESR spectra of probe measured as a function of temperature in semi IPN based on PI and PMMA with 50/50 composition [81].

mechanical properties up to a certain extent; however, higher crosslinking is not desirable. The combination of 2% crosslinker concentration in PI and 4% in PMMA in PI/PMMA IPN led to the development of nanostructured morphology and the said IPN performed better in mechanical measurements. For a given composition, the elongation at break decreased with the increase of crosslinking in the PI phase. This was due to the reduced chain flexibility of the PI matrix with higher crosslinking. Also crosslinking increased the PMMA entanglement density and interpenetration between the phases. The PMMA network density was found to have only little influence on the elongation at break of full IPNs. Young’s modulus was also found to increase with the increase in crosslinking of the rubber phase. Also crosslink density of the PMMA phase was found to have a pronounced effect in the Young’s modulus of full IPNs. The 6% crosslinker concentration in PMMA showed the highest value among the studied samples. It was found that shore A hardness increased with the rise in PMMA content. Moreover, main chain and segmental dynamics of PI and PMMA chains in semi IPNs with different composition, crosslink density, and molecular weight were studied over a wide range of temperatures, using the ESR spin probe technique and dynamic mechanical analysis and confirmed the dramatic changes in the motional behavior of both polymers due to the molecular level interpenetration between two polymer chains in synthesised semi IPN [82]. At lower temperatures, the ESR spectra approached the rigid limit spectrum, and as the temperature increased, the spectral lines got narrower and outer peaks shifted inward. In the range of 293 K-373 K, complex spectra were obtained, an indication of the presence of motionally distinct regions in semi IPNs (Figure 7.41). 7.3.1

Modeling of Mechanical Properties of IPNS

Further development and optimization of different mechanical properties of IPNs requires that experimental results be complemented by theoretical studies and computer simulations as they have the advantage of providing insights into the role of individual factors, and to rationally elucidate their significance.Various

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Tensile strength, MPa

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Experimental Parallel Series Helpsin-Tai Takayanaki Kunori 1 Kunori 2

Figure 7.42

x x x

x

x 0

10

20 30 40 wt-% of PMMA

50

60

Model fitting–tensile strength of IPNs based on PI/PMMA [81].

models, such as the parallel model, the series model, the Halpin–Tsai equation and the Takayanaki model were used to study the mechanical behavior of the IPNs. Kunori and Geil models were used to determine fracture paths (matrix or interface). John et al. applied the Kunori and Geil models to IPNs based on PI/PMMA. [81]. Figures 7.42 and 7.43 show the theoretical and experimental curves of the tensile strength and Young’s modulus values, respectively, 900 Experimental Parallel Series Halpsin-Tai Takayanaki Kunori 1 Kunori 2

800 700 600 Modulus, MPa

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Figure 7.43

10

20 30 40 Wt-% of PMMA

50

60

Model fitting–Young’s modulus IPNs based on PI/PMMA [81].

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1010

E′ (DYNE/CM2) AT 25 °C

RNP

109

108

107 0 PEAB

PINR

NORM. IPN INV. IPN BLEND GRAFT (27) 50 WT%

100 PS

Figure 7.44 Storage moduli at 25◦ C for PEAB-PS IPNs. R stands for rubber and P stands for plastic. The dotted line indicates the theoretical value of the storage modulus for two continuous phases, based on the Bauer model. Reprinted from [80]. Copyright (1972) with permission from American Chemical Society.

for the six models. For both tensile strength and modulus, parallel model fits more closely to the experimental values up to 50 wt% of PMMA and then the values deviate. This may be because both the PI and PMMA phases have a co-continuous morphology above 55 wt% PMMA, rather than having a matrix-dispersed phase morphology. The theoretical curve of Kunori (equation for fracture through matrix) comes closest to the experimental curve compared to the other models. Therefore, it may be concluded that the fracture path is through matrix rather than through the interface. The Kunori values superimpose over parallel model, so lines cannot be distinguished in the figures. The modulus-composition behavior of the PEA/S-co-MMA IPNs was analyzed with some of the theoretical equations based on mechanical models [80]. Employing the theory of Bauer et al., which is essentially a modification of the earlier mechanical models of Takayanagi, they judged the relative continuity of the two phases [83]. Figures 7.44 and 7.45 present storage modulus data as functions of composition for PEAB-PS and PEAB-PMMA IPNs, respectively. Also included are the solution blended material, a graft copolymer and data from a mixed latex. The upper solid lines represent a continuous plastic phase, while the lower solid lines represent a continuous elastomer phase. The dotted lines indicate the result obtained for two equally continuous phases, all according to theory. Allowing for the relative oversimplification of the mechanical model, the results tend to indicate that both phases do in fact exhibit some degree of continuity, especially the PEA-PMMA compositions. A possible exception is the 75/25 PEAB-PS point in Figure 7.44, which suggests a continuous elastomeric phase. It should be noted, however, that the theoretical curves were derived for uncrosslinked polyblends.

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E’ (DYNE/CM2) AT 25 °C

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108

107 0 PEAB

PINR

NORM. IPN INV. IPN BLEND (28) MIX LATX (25) CHCI3 (26) GRAFT (26) 50 WT%

100 PMMA

Figure 7.45 Storage modulus at 25” for PEAB-PMMA IPN’s.Within the limits of the Bauer theory, all such blends and IPNs appear to have two continuous phases. Reprinted from [80]. Copyright (1972) with permission from American Chemical Society.

7.4

Polymer Gels

A polymer gel is a crosslinked polymer network swollen in a liquid medium. Its properties depend strongly on the interaction of these two components. Polymer hydrogels can be divided into two main classes, i.e., chemically crosslinked hydrogels, which are composed of polymer networks with covalent bonding, and physically crosslinked hydrogels, which are composed of physical networks with noncovalent interactions. To date, hydrogels of poly(N-alkylacrylamides) such as poly(N-isopropylacrylamide) (PNIPA) and poly(N.Ndimethylacrylamide) (PDMAA), which have attracted extensive attention as water-absorbing, soft and stimulisensitive materials, were all prepared by chemical crosslinking reactions using an organic crosslinker such as methylene bis (acrylamide) (BIS) [84–99]. However. because of the random nature of the crosslinking reactions produced by a large number of organic crosslinkers, the conventional chemically crosslinked hydrogels (hereinafter abbreviated as OR gels) have many limitations in morphology and properties, e.g., morphological inhomogeneity, mechanical weakness, limited swelling at equilibrium, and slow response to stimuli. Therefore, for example, NIPA-OR gels consisting of PNIPA chemical networks readily become turbid due to structural inhomogeneities induced by increasing the crosslink density, pressure, and polymerization temperature [100–102]. Also, swelling ratios or de-swelling rates are not always sufficient for certain applications [88, 103, 104]. These limitations mainly

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exfoliated clay Dic g2

g1

χ

100 nm

Figure 7.46 Schematic representation of the structural model with organic/inorganic networks in the NC gel. Dic is interparticle distance of exfoliated clay sheets; χ , g1 and g2 represent crosslinked chain, grafted chain, and looped chain. Reprinted from [107]. Copyright (2002) with permission from John Wiley & Sons.

arise from restricted molecular motions of PNIPA chains caused by crosslinking with a large number of organic crosslinkers. Furthermore, the most serious limitations of OR gels are due to their weak and brittle nature [105, 106]. Irrespective of the composition or conditions of preparation, NIPA-OR gels always broke at elongations less than 30% or when bent through 180◦ [107] and so were very difficult to handle in applications requiring significant applied stress or strain. This limitation is due to the low average and broad distribution of chain lengths between crosslinking points in OR gels. Recently, and in order to overcome the limitations of OR gels, a new type of polymer hydrogel nanocomposite type hydrogels has been prepared (herein after abbreviated as NC gels, e.g., NIPA-NC gel) consisting of a unique organic polymer (e.g. PNIPA)/inorganic (clay) network (Figure 7.46) [107]. NC gels showed remarkable improvements in mechanical, optical, and swelling-deswelling properties. That is, NC gels simultaneously exhibit high transparency (structural homogeneity) irrespective of the preparation conditions, excellent tough mechanical properties, with astonishingly large elongations, large swelling ratios, and rapid deswelling responses to temperature changes. NIPA-NC gels were prepared by in situ free-radical

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polymerization of N-isopropylacrylamide, in the presence of inorganic clay exfoliated in an aqueous media. The formation of effective polymer/clay networks could be achieved using inorganic clays which act as multifunctional crosslinking agents through noncovalent interactions instead of using organic crosslinkers. The characteristics of NC gels strongly depend on their compositions. In other words, they could be controlled over a wide range by altering their composition and network structure. Haraguchi et al. prepared nanocomposite hydrogels composed of poly(N,N-dimethylacrylamide) (PDMAA) and clay by in situ free-radical polymerization of N,N-dimethylacrylamide (DMAA) in the presence of inorganic clay in aqueous solution [108]. DMAA·NC gels consist of organic/inorganic networks similar to those of NIPA-NC gels established previously [107]. The proposed model structure shown in Figure 7.46 consists of networks of inorganic clay crosslinked by PDMAA through noncovalent interactions (probably hydrogen bonding ionic and/or coordinate interactions) between the PDMAA chains and the clay. Here, it is noted that, in the NC gels, free linear PDMAA chains could not be detected during the purification and swelling processes. This suggests that all polymer chains are attached to the clay and involved in the network. Also, the PDMAA component of the NC gel forms not only crosslinked chains (x) but also grafted chains with a free chain end (g1 ) or looped chains (g2 ) in which both ends are attached to one clay sheet. Further, as for crosslinked chains, topologically crosslinked chains consisting of looped chains or trapped interchain entanglements, etc., might be included in NC gels, although the precise nature of the interactions and the mechanism of network formation are still under examination. The interactions between (P)DMAA initiator, and clay in the aqueous media and in the dried state were observed by similar means to those used for the PNIPA-c1ay system: viz. by solution viscometry and by IR spectroscopy. The proposed model for the nature of PDMAA/clay networks seems reasonable on the basis of analyses of their characteristic rubbery mechanical properties and their swelling behavior. To verify the proposed model, the clay dispersion and the chain flexibility of the polymer in DMAA-NC gels were examined. For the clay, the fact that the resulting NC gel and its dried gel were both optically transparent, almost regardless of clay content (Cclay ), indicates that it is finely and uniformly dispersed in the hydrogels. The nature of the clay dispersion was also elucidated by SAXS measurements. It was generally observed that DMAA-NC gels exhibited excellent tensile mechanical properties compared with conventional DMAA-OR gels. Similar to NIPA-OR gels, DMAA-OR gels broke at very low extension, ca. 50%, irrespective of crosslinker content (CBIS = 1 − 9) and showed very low ultimate tensile strengths (ca.7kPa). On the contrary, DMAA-NC gels generally exhibited very large elongations at break and high strengths. Also, it was observed that the tensile properties of DMAA-NC gels strongly depend on Cclay . Figure 7.47 shows tensile stress–strain curves measured for DMAA-NC gels with different Cclay , where both modulus and strength exhibit remarkable increases with increasing Cclay . On the other hand, elongations at break decreased slightly with increasing Cclay , particularly at relatively low Cclay (< NC4). In samples with higher Cclay (≥ NC4), the elongation at break changed little, staying at an almost constant value of ca. 1300%. The effects of clay content (Cclay ) on the mechanical properties of DMAA-NC gels described above were very similar to those for NIPA-NC gels [109], although the absolute values of properties and their detailed dependencies on composition were somewhat different. Changes in tensile modulus, tensile strength, and elongation at break by altering Cclay for both DMAA-NC-Ml and NIPA-NC-MI gels are summarized in Figure 7.48a–c. It was shown that DMAA-NC gels exhibit, as a whole, lower moduli and higher elongations at break than those of NIPA-NC gels, while the ultimate strengths were similar for both gels. The stress–strain behavior for a series of DMAA-NC2.5 and –NC5.5 gels with different polymer content (Cp ), respectively were studied. Figure 7.49 shows the stress–strain curves for a series of DMAA-NC2.5 with different (Cp ). Cp was varied over about 2 orders of magnitude, from Cp = 0.1 to 8. NC gels prepared using fixed conditions, apart from Cp , showed a large change in their mechanical characteristics, from very brittle to very tough, on increasing Cp . In the very low Cp region, a lower critical polymer content (C *p ), below which

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NC7

250

Stress [kPa]

200 NC5.5

150

NC4

100

NC3

50 NC2.5 0

0

500

1000 Strain (%)

1500

2000

Figure 7.47 Stress–strain curves of DMAA-NC-Ml gels with different clay content (NC2.S to NC7). All hydrogels had the same polymer/water ratio (= 1/10 (w/w). Reprinted from [108]. Copyright (2003) with permission from American Chemical Society.

NC gels became brittle, was observed at ca. 0.13 for both DMAA-NC2.5 and –NC5.5 gels. Hydrogels with Cp < C *p were too brittle to carry out standard tensile testing. Also, hydrogels prepared with Cp around C *p were often neither uniform nor transparent. In the relatively low Cp region (C *p ∼ ca.0.5), the mechanical properties of NC gels changed dramatically with increasing Cp . In particular, elongations at break increased markedly from near zero to 1100–1600%. This indicates that the formation of the fundamental network of NC gels requires this magnitude of Cp . Then, in the following intermediate Cp region, it was estimated that increasing Cp may result in an increase in the number of effective crosslinks. Actually, it was observed that the modulus and the strength clearly increased significantly with increasing Cp in this region (Cp = 0.5 – 5.5). On the other hand, the elongation at break was almost constant or changed little on increasing Cp . For example, in a series of NC2.5 gels, the elongations at break were observed to be constant at ca. 1600%, irrespective of Cp . However, in a series of NC5.5 gels, elongations at break gradually changed Cp in this region. Also, the absolute maximum elongations of NC5.5 gels were, on the whole, a little less than those of NC2.5 gels. Mechanical strength of polymer gel electrolytes was investigated on the basis of polymer-solvent affinity [110]. Four different polymers – poly(vinylidene fluoride) (PVdF), poly(vinylidene fluoride) (PVdF)hexafluoropropylene (HFP) copolymer, poly(acrylonitrile) (PAN) and poly(methyl methacrylate) (PMMA) – were employed as the gel-forming polymer matrix, while ethylene carbonate (EC):propylene carbonate (PC) was used as the plasticizing solvent. The solvent retention ability of polymer gels decreases in the order of PMMA ≥ PAN  P (VdF-HFP) ≥PVdF, which is surely the relative order of polymer affinity for the solvent. In Figure 7.50, the elastic modulus (slope of stress/strain curve in the elastic region), ultimate tensile strength (stress at break) and ultimate elongation (elongation at break) were compared between the polymer gel films. P(VdF-HFP) gel shows a higher elastic modulus and tensile strength as compared to the PMMA- or PAN-based gels. It can thus be generalized that polymer gels are mechanically stronger when polymer matrix has a low affinity for solvent molecules, in which circumstances a microscopic phase transition is frequently observed.

7.5 Polymer Composites During the last three decades, composites have replaced traditional materials in many engineering applications due to their excellent specific properties. Fillers play important roles in modifying the desirable properties of

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DMAA-NC4-M1*

NIPA-NC-M1

35

Modulus [kPa]

30 25 20 15 10 5 0 (a) Modulus

Strength [kPa]

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Elongation [%]

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0

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4 Cclay

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(c) Elongation at break

Figure 7.48 Changes of (a) tensile modulus, (b) tensile strength, and (c) elongation at break by altering clay contents for two series of NC gels: DMAA-NC-MI gels (closed circle: solid line) and NIPA-NC-MI gels (open square: dotted line). The properties of DMAA-NC4-MI* gel with the same polymer weight with NIPA-NC4-MI gels are also plotted. Reprinted from [109]. Copyright (2003) with permission from American Chemical Society.

polymers and reducing the cost of their composites. The fibers such as glass fibers, aramide fibers, carbon fibers and natural fibers have been used as macrofillers to develop thermoplastic and thermoset composites. Inorganic fillers with dimensions in the micrometer range, e.g., calcium carbonate, glass beads, silica, carbon black and talc, have been used extensively to enhance the mechanical properties of polymers. Such properties can indeed be tailored by changing the volume fraction, shape, and size of the filler particles. It is logical to anticipate that the dispersion of fillers with dimensions in the nanometer level having very large aspect ratio

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250 Cp = 5.5

DMAA – NC2.5 gel Cp = 8

200 Stress [kPa]

Cp = 4.5

150 Cp = 2

100

50 Cp = 0.13

0

0

500

Cp = 1

Cp = 0.22

Cp = 0.5

1000 Strain (%)

1500

2000

Figure 7.49 Effect of polymer contents (Cp = 0.13 − 8) on the stress–strain curves of DMAA-NC2.5 gels. All hydrogels had the same clay/water ratio (= 0.23/10 (w/w). Reprinted from [109]. Copyright (2003) with permission from American Chemical Society.

and stiffness in a polymer matrix could lead to even higher mechanical performances. These fillers include layered silicates, nanosilica and carbon nanotubes. 7.5.1

Mechanical properties of polymer macrocomposites

In the field of composites, the fiber reinforcement of matrices was initially developed using man-made fibers such as glass, carbon, aramid, etc., in order to take advantage of their high tensile moduli. Polymer-matrix composites, such as carbon or glass-fiber reinforced plastics (CFRP/GFRP) have been widely used in industry since they have high strength and modulus. In fact, the total amount of consumption of GFRP was about 382 thousand tons in Japan in 2001. 2.0 P(VdF-HFP) gel

Stress/MPa

1.5

1.0

0.5 PAN gel 0.0

0

100

200 Elongation/%

PMMA gel 300

Figure 7.50 Stress–strain curves of the polymer gel films. Reprinted from [110]. Copyright (2001) with permission from Elsevier.

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40

SBSS (MPa)

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30

25

20

15

0

5

10

15 20 25 30 Cooling rate (°C/min)

35

40

Figure 7.51 Measured short beam shear stresses (SBSS) on the PP/G and m-PP/s-G composites as a function of CR. Reprinted from [111]. Copyright (1998) with permission from John Wiley & Sons.

Youssef and Denault investigated the effect of the interfacial improvement by the addition of a chemically modified PP and a specific glass fiber thermoplastic sizing on the mechanical properties of the composites [111]. Mechanical tests were carried out on both PP/G (consisting of a pure PP matrix reinforced with 52 wt% of E-glass fibers) and m-PP/s-G (consisting of a blend of chemically modified PP with pure PP reinforced with 60 wt% of E-glass fibers coated with a specific thermoplastic sizing) composites. The study of the interface quality via the measurement of the short beam stresses and the response of the composite under tension in the ±45◦ direction showed that the presence of the modified PP in the pure PP matrix and the fiber sizing have noticeable effects on the interfacial strength. In fact, the m-PP/s-G system showed higher SBS (Short Beam Shear) stress values (Figure 7.51), higher tensile moduli and higher failure stresses in the ±45◦ direction. However, these performance parameters are significantly reduced when the composites are cooled down very slowly resulting in large spherulitical crystalline structures. This suggests that the contribution of the amorphous phase to the fiber-matrix interaction in this composites is important. When the thermosetting resins are used as a matrix for glass fiber, it is usually difficult to recycle the material. The thermoplastics are also impractical to recycle due to high cost and low quality issues in the present state. Most wasted FRPs are dumped although they do not decompose itself naturally in the ground, while others are burned. In recent years, a serious problem has come up in the use of plastics, especially for the polymer composite materials. Energy recycling systems are also under development using polymer composites as solid fuel. However, the glass fibers in the composites reduce the net heat and might damage the furnace, if their composites are burned in it as a solid fuel. It introduces another problems such as the disposal of the remains, because the glass fibers in the composite would also remain in the incinerator. Meanwhile, natural fiber composites are claimed to offer environmental advantages such as reduced dependence on nonrenewable energy/material sources, lower pollutant emissions, lower greenhouse gas emissions, enhanced energy recovery, and end-of-life biodegradability of components. Such superior environmental performance is an important driver of increased future use of natural fiber composites, Tensile, flexural, and impact behavior of PALF-reinforced polyester composites as a function of fiber loading, fiber length, and fiber surface modification were investigated by Devi et al. [112]. The PALF–polyester

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5

5 FLEXURAL STRENGTH

FLEXURAL STRENGTH

10 mm

4

70

3

60

50 (a)

2

0

10

20

30

FIBRE LENGTH, mm

40

1 50

FLEXURAL STRENGTH, (MPa)

FLEXURAL STRENGTH, MPa

FLEXURAL MODULUS

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FLEXURAL MODULUS

30 mm

80 4 60 3 40 2 20

0 (b)

297

0

5

FLEXURAL MODULUS (GPa)

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Figure 7.52 (a) Variation of flexural strength and flexural modulus with fiber length (fiber loading 30 wt%); (b) Variation of flexural strength and flexural modulus with fiber loading of PALF–polyester composites [112].

composites exhibit superior mechanical properties when compared to other natural-fiber polyester composites and can be used as structural composites. The optimum length of the fiber required to obtain PALF–polyester composites of maximum properties was found to be 30 mm (Figure 7.52a). The stress–strain behavior in tension reveals that neat polyester is brittle and the addition of fibers makes the matrix more ductile. The tensile strength and Young’s modulus of PALF polyester composites increased linearly with the fiber weight fraction. But in the case of flexural strength, there is a leveling off beyond 30% (Figure 7.52b). The impact strength also increased linearly with the weight fraction of the fiber. The composite with 30 wt% fiber content exhibits an impact strength of 24 kJ/m2 . The high toughness of this natural fiber polymer composite places it in the category of tough engineering materials. A significant increase in the strength of the composites was observed after treatment of the fibers. The best improvement was observed in the case of silane A-172-treated fiber composites. Okubo et al. developed composites for ecological purposes (Eco-composites) using bamboo fibers extracted by the steam explosion technique [113]. The bamboo fibers (bundles) had a sufficient specific strength, which is equivalent to that of conventional glass fibers. The tensile strength and modulus of PP-based composites using steam-exploded fibers increased about 15 and 30%, respectively, due to well impregnation and the reduction of the number of voids, compared to the composite using fibers that are mechanically extracted. The steam explosion technique is found to be an effective method to extract bamboo fibers for reinforcing thermoplastics The influence of various chemical treatments on the properties of sisal/PE composites has been investigated by Joseph et al. [114]. The chemical treatments included treatments with sodium hydroxide, isocyanate, and peroxide. The enhancement in the properties was ascribed to the bonding between sisal fiber and the PE matrix. Treatment with the cardanol derivative of toluene isocyanate was found to be better than other treatments as evidenced by the decrease in the hydrophilic nature of the composite. The composites exhibited better dimensional stability and retention of properties even after aging, which was ascribed to the improved moisture resistance.

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7.5.2

Mechanical Properties of Polymer Microcomposites

There is now considerable evidence that microfillers can significantly affect the structure of the matrix polymer itself, and hence the properties of the final composite. Commonly-used micro -fillers for polymers include mineral fillers such as talc, calcium carbonate, crystalline silica, and synthetic fillers such as carbon black, synthetic silica, etc. The toughening of polypropylene with rigid particles leads to a system with higher stiffness and higher impact resistance. A polypropylene–CaCO3 composite was processed by Zuiderduin et al. [115] which had a significant higher modulus and simultaneously showed improved toughness. The notched Izod impact energy could be raised from 2 to 50– 60 kJ/m2 at room temperature while increasing the modulus. Effect of processing conditions on the dynamic mechanical behavior of carbon black (CB) filled ethylene/ethylacrylate copolymer (EEA) composites was investigated [116]. The compounds were prepared by two methods, solution blending and mechanical mixing. Compared with the solution counterpart, the mechanical composites have a high dynamic elastic modulus which results from the good dispersion state of carbon black in EEA, i.e., the strong interaction between carbon black and EEA. 7.5.3

Mechanical Properties of Polymer Nanocomposites

Polymeric nanocomposites can be considered as an important category of organic–inorganic hybrid materials in which inorganic nanoscale building blocks (e.g., nanoparticles, nanotubes, or nanometer-thick sheets) are dispersed in an organic polymer matrix [117–121]. When compared to conventional composites based on micrometer-sized fillers, the interface between the filler particles and the matrix in polymer nanocomposites constitutes a much greater area within the bulk material, and hence influences the composite’s properties to a much greater extent, even at a rather low filler loading [122–124]. Polymer nanocomposites reinforced by relatively small amounts of ultra-fine, nano-particles (most often clay platelets) proved exceptionally promising engineering materials with an unexpectedly high stiffness/toughness ratio, gas barrier properties, flame retardence, etc. The real interest in nanotechnology is to create revolutionary properties and functions by tailoring materials and designing devices on the nanometer scale. According to a report, the total worldwide market for polymer nanocomposites reached 11.1 million Kg valued to US$90.8 million in 2003. This market was expected to grow at an average annual growth rate of 18.4% to reach US$211 million by 2008. 7.5.3.1

Factors Affecting Mechanical Properties of Nanocomposites

Influence of Filler Dispersion on Mechanical Properties Mechanical properties of polymer–clay nanocomposites are highly related to their microstructure which in turn is directly related to the exfoliation and dispersion of clay platelets in the polymer matrix. Tensile properties of the MA compatibilized PE–organoMMT nanocomposites were studied by Lee et al [125]. The tensile strength and modulus tend to increase with increasing clay content (Figure 7.53a). Such an increasing trend is more obvious for the tensile modulus. The increase in the tensile strength is higher at low clay content, indicating that the clay layers are better exfoliated. The reinforcing effect is lower for nanocomposites with higher clay content owing to some clay platelets being partially exfoliated and stacked. From Figure 7.53b, the strain at break decreases with increasing clay content as expected. The complete exfoliated morphology of polymer-MMT nanocomposites contributes greatly to their impact property and modulus, while the intercalated state of partially exfoliated state contributes more to the final material’s modulus. In the case of PP-based clay nanocomposite, how the dispersion state of MMTs decides the final properties of izod impact properties is shown in Figure 7.54 [126].

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Figure 7.53 (a) Tensile modulus and strength and (b) strain at break vs. clay content for the MAcompatibilized PE–organoMMT nanocomposites. Reprinted from [125]. Copyright (2005) with permission from Elsevier.

Liu et al. synthesized organoclay-modified epoxy nanocomposites by high pressure mixing method (HPMM) [116]. Better dispersion of organoclay in epoxy matrix has been achieved by HPMM than direct mixing method (DMM) (Figure 7.55). The nanocomposites formed by the HPMM showed a dramatic improvement in fracture toughness at very low clay loading; that is, K1C and G1C were increased by 1.7 and 3.2 times respectively, at 1.5 phr (about 1 wt%) organoclay loading (Figure 7.56).

44 Exfoliated state

Izod Impact Strength (J/m)

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36

32 Blending Intercalated state 28

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Figure 7.54 Relationship between the izod impact property–clay load in polypropylene-based nanocomposites. Reprinted from [126]. Copyright (2005) Elsevier.

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Figure 7.55 SEM micrographs of fracture surfaces of nanocomposites made with (a) DMM showing agglomerates clay layers in epoxy matrix; (b) HPMM showing fine dispersion of clay layers in the epoxy matrix. Reprinted from [127]. Copyright (2005) with permission from Elsevier.

Influence of Filler Content on the Mechanical Properties The amount of filler strongly influences the mechanical properties of polymer nanocomposites. Zhou et al. studied the effect of filler content on the mechanical properties of multi-walled carbon nanotubes (CNTs) reinforced epoxy [128]. When dispersing CNT in polymer matrix, it is important to keep the filler volume (or weight) fraction below a certain value to maintain the strength and fracture toughness. Optimal loading of CNT in matrix is a key parameter to developing multifunctional nanophased composite. Flexural modulus steadily increases with a higher CNT weight percent. Modulus improved by 11.7% with an addition of a 0.4 wt% of CNTs. Flexural strength and fracture toughness peaked in a 0.3 wt% CNT/epoxy system. The decrease in strength and fracture toughness in 0.4% CNT/epoxy was attributed to poor dispersions of nanotubes in the composite. Compared to neat epoxy, 500

2 1.30E with HPMM

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Figure 7.56 (a) Critical strain energy release rate (G1C ); (b) Critical stress intensity factor (K1C ) of nanocomposites and filler composites [127].

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Table 7.6 Mechanical properties of neat and CNT modified epoxy [128]. Modulus [GPa] Neat epoxy 0.1% CNT 0.2% CNT 0.3% CNT 0.4% CNT

2.46 2.54 2.60 2.65 2.75

Strength [MPa] 93.5 109 115 121 113

Failure strain [%] 4.02 6.06 6.80 7.58 5.12

DMA results indicated a 93% improvement in storage modulus in 0.4 wt% CNT/epoxy at room temperature and a 17◦ C increase in Tg (Table 7.6, Figure 7.57). The approach for incorporating nanoscopic inorganic cluster into organic polymer is to design well-defined inorganic oligomers with a single polymerization site per cluster. Each oligomeric cluster has an exactly defined degree of polymerization of eight, (RSiO1.5 )8 , or more precisely, P1 R7 Si8 O12 , where R and P are organic groups. These polyhedral oligomeric silsesquioxane (POSS) macromers have an inorganic silica-like core and are surrounded by eight organic groups, of which seven are inert and just one is reactive. Further polymerization involving the single reactive P site results in a linear polymer containing monodisperse, nanosize inorganic clusters pendent to an organic polymer backbone. These hybrid inorganic–organic polymers can be processed further like any thermoplastic polymers. Lee and Lichtenhan have modified epoxy with polymerizable polyhedral oligomeric silsesquioxane (POSS) macromers [129]. Figure 7.58, depicts the relaxation modulus curves tested for different loading of POSS–epoxy monomers after 64 h of isothermal aging at a temperature of 63.9◦ C. It is clear that these relaxation curves can be superimposed with only horizontal shifts along the time axis. However, it is interesting to point out that the value of E0 is not affected by the presence of the nanoreinforcement. This may be in part a reflection of the monofunctional nature in POSS epoxy

Fracture toughness [MPa.mm1/2]

180

160

140

120

100 0.00

0.10 0.20 0.30 Weight fraction of CNTs [%]

0.40

Figure 7.57 Effect of CNT contents on fracture toughness of epoxy. Reprinted from [128]. Copyright (2008) with permission from eXPRESS Polymer Letters.

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ta=64 hours

101 T=63.9°C ε=0.1%

Relaxation Modulus (GPa)

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Time (Seconds)

Figure 7.58 Small-strain stress relaxation modulus curves for DGEBA–D230–Heloxy 67–wt % POSS–epoxy glass after 64 h of isothermal aging at a temperature of 63.9◦ C. Applied strain is 0.001[129].

monomers used in this study. Despite the ability of the POSS cages to hinder the relaxation motion of network junctions from a chain terminus location within the network, they do not contribute to the overall deformation process of such glassy networks from this position. Interestingly, such monofunctional POSS–epoxide may be useful for enhancing glass transition without increasing crosslink density and potentially detracting from the desirable mechanical properties of such epoxy networks. Filler-Matrix Adhesion The effect of clay modification on organo-MMT/NBR nanocomposites was studied by Kim et al. [130]. They modified the organoclays with alkylamine cations. Alkylamines used were octylamine (C8), DDA (C12), and ODA (C18). The organo-MMT content in the prepared nanocomposites was fixed at 0.0, 1.86, 4.52, 8.70, 12.45, and 15.94 wt%. For the C8-MMT/NBR nanocomposites, the tensile strength increased rapidly with increasing clay content from 0 to 4.52 wt%, but the change was less when the clay content increased beyond 4.52 wt%. In the case of C12-MMT and C18-MMT, the tensile strength increased rapidly with the clay content up to 8.7 wt% beyond which there was very little change. For the C8-MMT/NBR system, the tensile modulus increased slightly with increasing clay content. But the tensile modulus of C12-MMT/NBR and C18-MMT/NBR nanocomposites increased rapidly with increasing clay content. Below 8.7 wt%, the improvement in elongation at break may be attributed in part to the plasticizing effect of alkyl ammonium ions that are located at the clay–NBR interface. Above 8.7 wt%, the lowering of the elongation at break was explained by the formation of nonexfoliated aggregates at higher clay content that made these composites much stiffer. The differences in mechanical properties among the C8-MMT/NBR, C12-MMT/NBR, and C18-MMT/NBR hybrids were explained by the differences in hydrophobicity of the

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Flexural strength (Mpa)

Mechanical and Viscoelastic Characterization of Multiphase Polymer Systems 135 130 125 120 115 110 105 100

E+OC series E+UC series 0

2 4 6 8 10 clay content (wt%)

Figure 7.59

Flexure modulus (GPa)

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4.5 4 3.5 3 E+OC series E+UC series

2.5 2 0

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(a) Flexural strength of epoxy/clay series; (b) Flexural modulus of epoxy/clay series [131].

organo-MMT. Overall, the mechanical properties increased in the order C8-MMT < C12-MMT < C18-MMT, depending on the length of the alkyl chain in the alkyl ammonium. The diglycidyl ether of bisphenol A (DGEBA) epoxy resin system filled with organo clay (OC) and unmodified clay (UC) were processed and Mohan et al. studied the effect of filler-matrix adhesion on mechanical properties [131]. The flexural properties of the epoxy filled with organo and unmodified clay particles are shown in Figure 7.59. The organoclay-filled epoxy has better improvement in the flexural modulus than that of the unmodified clay-filled epoxy. The important parameter that affects this property by incorporating such fillers is the quality of interface in the composites, i.e., the adhesive strength and the interfacial stiffness of the composite medium. These two factors play a crucial role in stress transfer and the elastic deformation from the matrix to the fillers. This is very much applicable to the nanoparticle-filled polymers, due to high surface area of the organoclay filler which increases the contact area to the matrix and imparts a high portion of interface. If the interface is poor between the matrix and fillers, there is much less chance for the fillers to carry the load and results in less modulus, which is seen for the unmodified clay-filled epoxy polymer. As a result of this, the unmodified clay-filled composite cannot have good strength as compared to the matrix material. But in the nanocomposites, because of the enhanced interfacial property owing to their large surface area of fillers, the increased property is observed which reveals that stresses are effectively transferred through the interface. For polymer-based nanocomposites, an appropriate surface treatment of inorganic nanoparticles should not only improve dispersion of the fillers, but also bring about notable influence on the interfacial characteristics, and subsequently enhance the mechanical properties of the ultimate composites. The internal surfaces (interfaces) are critical in determining the properties of polymer silica nanocomposites. Taking the advantage of graft polymerization, Cai et al. [132] prepared nanosilica-modified polypropylene with specific structural requirements in one step. The key issue lies in the introduction of a polymeric foaming agent containing soft segments (i.e., poly(p-vinylphenylsulfonylhydrazide-co-butyl acrylate)) onto the surface of nano-SiO2 . In the case of melt blending with PP, the side sulfonyl hydrazide groups on the grafted copolymer are able to be gasified like foaming agent to induce a localized bubble-stretching effect that pulls apart nanoparticles agglomerates. Meanwhile, the skeleton of the grafted copolymer would get entangled with the matrix polymer forming strong interfacial interaction, and the poly (butyl acrylate) units in the grafted copolymer help to raise ductility of the interlayer (Figure 7.60). Compared to the composite containing rigid macromolecular foaming agent grafted nano-SiO2 (i.e., poly (p-vinylphenylsulfonylhydrazide) grafted nano-SiO2/PP), the composite fabricated in this work (i.e., poly (pvinylphenylsulfonylhydrazide-co-butyl acrylate) grafted nano-SiO2/PP) shows much greater increment in notched impact strength without expense of lowing tensile performance.

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Handbook of Multiphase Polymer Systems Residual sulfonyl hydrazide group PBA chain units O

OH

OCH2 Si OCH2

PP matrix Graft copolymerization

(CH2)3

In situ bubble stretching

O=O

C CH2 CH

SO2NHNH2

n

SiO2 nanoparticles

7.5.4

CH2 CH C O C4H9 m

Figure 7.60

Bubble generated by gasification of side sulfonyl hydrazide groups on the grafted polymer Bubble-stretching induced dispersion of SiO2 nanoparticles in PP matrix and simultaneous interfacial tailoring

O

Schematic drawing of the proposed route for making nano-SiO2/PP composites [60].

Mechanical Modeling of Polymer Nanocomposites

Several models have been developed to predict the mechanical behavior of polymer nanocomposites, including Halpin–Tsai, Mori–Tanaka, etc. The models generally include parameters such as the aspect ratio, volume fraction and the orientation of the reinforcement. The elastic modulus of the composite material can be predicted from the Halpin–Tsai equation assuming the fibers are discontinuous and aligned uniaxially [133, 134]. The longitudinal elastic modulus of composites Ec , is given by: Ec 1 + ξ ηφ = Em 1 − ηφ

(7.15)

where Em is the tensile modulus of the matrix and ϕ is the volume fraction of fiber reinforcement. The Mori–Tanaka mean field theory is used to assess the overall properties such as the effective stiffness tensor C* of the composites. It is based on the Elsheby method for estimating stress state in composite reinforced with misfitting inclusions. The composite is assumed to be composed of a continuous matrix and discrete of inclusions of different stiffness. The effective stiffness tensor C* is given by the following relation [135, 136]: C ∗ = C1 + V2 {(C2 − C1 )}

(7.16)

where C1 is the matrix phase stiffness tensor, C2 the inclusion stiffness tensor, V 2 the inclusion volume ratio, and A is the concentration tensor. For a composite consisting of a single, arbitrarily shaped inclusion perfectly bonded inside the matrix, the dilute strain concentration tensor of the effective particle is given by: A(di1) = [I + SC −1 (C2 − C1 )]−1

(7.17)

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where I is the fourth order unit tensor and S the fourth order Elshelby tensor. As the inclusion volume fraction increases, interaction between the inclusions reduces the accuracy of the dilute approximation.In other words, interactions of the field from other inclusions are expected to influence the evolution of the average fields in the matrix and the reinforcement. The Mori–Tanaka approach includes the effect of particle interaction. [137] In this case, A can be expressed as: A = A(di1) [V1 I + V2 {A(di1) }]−1

(7.18)

where V 1 is the matrix volume ratio. The Mori–Tanaka model has better predictive capability for fillers with relatively high aspect ratios. Tandon and Weng based on the Mori–Tanaka approach and derived the longitudinal modulus (E11 ) of the composite reinforced with platelets [138]: E 11 1 = Em 1 + φ f [−2νm A3 + (1 − νm )A4 + (1 + νm )A5 A]/2A

(7.19)

where ν m is the Poisson’s ratio of the matrix, and A, A3 , A4 and A5 are calculated from the matrix and filler properties and the components of the Elshelby tensor. In polymer–clay nanocomposites, parameters associated with hierarchical morphology of the clay such as the silicate interlayer spacing (d0 0 1 ), gallery spacing, platelet thickness, etc., should be incorporated into the micromechanics model. Brune and Bicerano [139] modified the Halpin–Tsai equation for tensile modulus of intercalated (or incompletely exfoliated) nanocomposites as: Ec 1 + ξ  η φ  = Em 1 − η φ 

(7.20)

where Er  is the ratio of the modulus of the platelet stack to that of the matrix, ξ  the aspect ratio of the platelet stack and η is the volume fraction of platelet stacks in the matrix. Recently, Fornes and Paul [140] emphasized the importance of aspect ratio and exfoliation ratio of clay in modeling stiffness. They performed simple calculations on the aspect ratio (l/t) of MMT platelets of the PA6–organoclay nanocomposites. They quantified those parameters and used the composite theories of Halpin–Tsai and Mori–Tanaka to predict the stiffness of high molecular weight nylon 6/clay nanocomposites and obtained good agreement with experimental data (Figure 7.61). They assumed that the composite consists of a matrix and stacks of clay platelet sheets. For the particle thickness (t) determination, they incorporated several parameters such as the silicate interlayer spacing (d0 0 1 ), number of platelets per particle and the thickness of an MMT platelet. From this, the number average particle thickness was determined to be 1.61 nm. On the basis of image analysis from TEM micrographs of nanocomposites, they determined the average particle length to be 91 nm. This resulted in an aspect ratio of 57. For an exfoliated structure, the platelets were completely delaminated and dispersed independently in the matrix with a thickness of 0.94 nm. Thus an aspect ratio of 97 was determined for an exfoliated structure. They substituted such particle aspect ratio values together with the stiffness of MMT (178 GPa) and PA6 (2.75 GPa) into the Halpin–Tsai equation. In Figure 7.61(a), experimental modulus data was compared with model predictions for aligned layered aluminosilicate nanocomposites having an aspect ratio identical to the experimental value of 57 (number average aspect ratio). The ratio of 97 corresponds to perfectly exfoliated morphology, i.e., the number average particle length, 91 nm, divided by the thickness of an individual platelet, 0.94 nm. It is clear from Figure 7.61(a) that Halpin–Tsai equations slightly overpredict the experimental data, while the Mori–Tanaka theory underpredicts the experimental data. However, both theories demonstrate that even higher levels of reinforcement might be possible with higher levels of exfoliation, larger platelet diameters, improved orientation, etc. Figures 7.61(b) and (c) compare the experimental data to model predictions based on an intercalated morphology (stacks of

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E/Em

Em = 2.75 GPa vf = 0.20 vm = 0.35

1/t 97

Halpin-Tsai

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Figure 7.61 Comparison of experimental and theoretical (based on unidirectional reinforcement) stiffness for high molecular weight nylon 6 nanocomposites. (a) Pure montmorillonite having a filler modulus of 178 GPa and aspect ratio of 57 (experimentally determined number average value) and 97, corresponding to complete exfoliation, and (b, c) intercalated morphology having one or more platelets per stack [129].

clay intercalated with polymer having one or more platelets per stack). The stack properties were based on experimental data, i.e., the stacks are 91 nm in length, have a repeat spacing of 1.8 nm, and each individual disk has a modulus of 178 GPa. It can be seen that the experimental data lie between the Halpin–Tsai curves corresponding to 1 and 2 platelets per stack, which is very close to the experimental determined value of 1.4. However, when Mori–Tanaka theory is used, the experimental data matches a completely exfoliated morphology, i.e., n = 1.

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7.6 Conclusion, Future Trends and Challenges The research in the area of multiphase polymer systems, especially polymer blends and composites, has got much attention for the past several decades in polymer science and technology. In this chapter, we have presented the various aspects concerning the mechanical and viscoelastic properties of multiphase polymer systems based on thermoplastics, rubbers and thermosets. The existing literature on mechanical properties of polymer blends including nanostuctured blends, polymer composites, IPNs and polymer gels have been discussed. The critical parameters which tune the mechanical properties of multiphase system are demonstrated with examples. The applicability of various micromechanical models on different multiphase polymer systems is included in this chapter. More effort needs to be placed in the field of computation modeling aiming at predicting the mechanical behavior of multiphase polymer systems. Detailed studies on the mechanical modeling of polymer gels are also required for predicting the mechanical properties. It should be emphasized that there is limited work on analytical modeling of the strength and toughness of polymer nanocomposites. There are several conflicting reports on the mechanical properties of polymer nanocomposites. More careful and in-depth analysis should be done to correlate the microstructure and mechanical properties of polymer nanocomposites. It is also important to study the effect of different mixers on the viscosity and hence on the morphology and properties of multiphase polymer systems. For the better marketing of the products we also need to look at the cost–performance ratio. Finally, it is important to add that a comprehensive understanding of microstructure and mechanical properties is very important for the fabrication of useful products from multiphase polymer systems.

References 1. L. A. Utracki, in Polymer Blends Handbook, L.A. Utracki (Ed.), Volume 1, Kluwer Academic Publishers, Dordrecht (2002). 2. I. W. Hamley, The Physics of Block Copolymers, Oxford, Oxford Science Publications (1998). 3. S. T. Milner, Macromolecules, 27, 2333 (1994). 4. H. Bramfeldt, P. Sarazin and P. Vermette, Polymer Degradation and Stability, 93, 877–882 (2008). 5. A. Kelarakis and K. Yoon, European Polymer Journal, 44, 3941–3945 (2008). 6. P. Tanpaiboonkul, W. Lerdwijitjarud, A. Sirivat and R.G. Larson, Polymer, 48, 3822–3835 (2007). 7. B. G. Sumpter, D. W. Noid and M. D. Barnes, Polymer, 44, 4389–4403 (2003). 8. C. Harrats, S. Thomas and G. Groeninckx, Micro and Nanostructured Multiphase Polymer Blend Systems: Phase Morphology and Interfaces, Taylor and Francis Group (2006). 9. D. S. Lee and S. C. Kim, Macromolecules, 17, 11 (1984). 10. G. M. Kavanagh, S. B. Ross-Murphy, Prog. Polym. Sci., 23, 533–562 (1998). 11. Y. Osada and J. P. Gong, Adv. Mater., 10, 11 (1998). 12. S. G. Advani, Processing and Properties of Nanocomposites, World Scientific Publishing Co. Pte. Ltd (2007). 13. D. R. Paul and C. B. Bucknall (Eds). Polymer Blends (Vol.1), Formulation (Vol.2), Performance, John Wiley & Sons, Inc, New York (2000). 14. C. W. Macoscko, Macromol.Symp., 149, 171–184 (2000). 15. H. Varghese, S. S. Bhagawan, S. S. Rao and S.Thomas, Eur. Polym. J., 31, 10, 957–967 (1995). 16. C. R. Kumar, K. E. George and S. Thomas, Journal of Applied Polymer Science, 61, 2383–2396 (1996). 17. H. Veenstra, P. C. J. Verkooijen, B. J. J. van Lent, J. van Dam, A. P. de Boer and A. P. H. J. Nijhof, Polymer, 41, 1817–1826 (2000). 18. Z. Oommen and S. Thomas, J. Appl. Polym. Sci., 65, 1245 (1997). 19. M. E. Broz, D. L. Vander Hart and N. R. Washburn, Biomaterials, 24, 4181–4190 (2003). 20. S. C. George, K. N. Ninan, G. Groeninckx and S. Thomas, Journal of Applied Polymer Science, 78, 1280–1303 (2000).

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21. T. Semba, K. Kitagawa, U. S. Ishiaku and H. Hamada, Journal of Applied Polymer Science, 101, 1816– 1825 (2006). 22. J. H. Chen, J. C. Zhong, Y. H. Cai, W. B. Su and Y. B. Yang, Polymer, 48, 2946–2957 (2007). 23. L. A. Utracki and R. Simha, Macromolecules, 37, 10123 (2004). 24. S. D. Burnside and E. P. Giannelis, Chem Mater., 7, 1596 (1995). 25. C. Creton, E. J. Kramer, H. R. Brown and C. Y. Hui, Adv. Polym. Sci., 156, 53 (2001). 26. S. Jose, PhD thesis, Mahatma Gandhi University (2006). 27. J. Jaafar, A. F. Ismail and T. Matsuura, Journal of Membrane Science, 345, 119–127 (2009) 28. G. Spadaro and G. Rizzo, Eur. Polym, J., 25, 1189 (1989). 29. S. J. Kim, S. R. Chowduhury, W. J. Cho and C. S. Ha, J. Appl. Polym. Sci., 89, 5, 1305 (2003). 30. N. Moussaif and G. Groeninckx, Polymer, 44, 7899 (2003). 31. G. Filippone, N. T. Dintcheva, D. Acierno and F. P. La Mantia, Polymer, 49, 1312 (2008). 32. M. Feng, F. L. Gong, C. Zhao, G. Chen, S. Zhang and M. Yang, Polym. Int., 53, 1529 (2004). 33. Y.J. Li and H. Shimizu, Macromol. Rapid Commun., 26, 710 (2005). 34. B. B. Khatua, D. J. Lee, H. Y. Kim and J. K. Kim, Macromolecules, 37, 2454 (2004). 35. F. C. Chiu, H. Z. Yen and C. E. Lee, Polymer Testing, 29, 397–406 (2010). 36. P. Motamedi and R. Bagher, Materials & Design, (2010). 37. A. Rahmatpour, M. Abdollahi and M. Shojaee, Journal of Macromolecular Science, Part B: Physics, 47, 523–531 (2008). 38. L. T. Vo and E. P. Giannelis, Macromolecules, 40, 8271–8276 (2007). 39. S. Joseph and S. Thomas, J. Polym. Sci. B: Polym. Phys., 40, 755–764 (2002). 40. R. C. Willemse, A. Speijer, A. E. Langeraar and A. Posthuma de Boer, Polymer, 40, 6645–6650 (1999). 41. A. Y. Coran and R. Patel, J. Appl Polym. Sci., 20, 3005 (1976). 42. T. Kunori and P. H. Geil, J Macromol Sci Phys B, 18, 135 (1980). 43. E. H. Kerner, Proc. Phys Soc., 696, 808, (1956). 44. A. H. J. Nijhof, LTM-report 1167, Delft University of Technology, The Netherlands, 1998. 45. J. J. Rajasekaran, J. G. Curro and J. D. Heneycutt, Macromolecules, 28, 6843 (1995). 46. M. Wang and N. Pan, Materials Science and Engineering R, 63, 1–30 (2008). 47. B. D. Favis, J. Appl. Polym.Sci., 39, 285 (1990). 48. J. Lu, Z. Qiu and W. Yang, Polymer, 48, 4196–4204 (2007). 49. J. M. Raj and C. Ranganathaiah, Polymer Degradation and Stability, 94, 397–403 (2009). 50. J. M. Huang and S. J. Yang, Polymer, 46, 8068–8078 (2005). 51. A. Robard, D. Patterson and G. Delmas, Macromolecules, 10, 706 (1977). 52. L. E. Fried, M. R. Manna and P. F. Pagoria, Annual Review of Materials Research, 31, 29, (2001). 53. S. Joseph, Ph.D thesis, Mahatma Gandhi University (2004). 54. K. A. Moly, S. S. Bhagawan, G. Groeninckx and S. Thomas, J. Appl. Polym. Sci., 100, 4526 (2006). 55. B. John, K. T. Varughes, Z. Oommen, P. P¨otschke and S. Thomas, Journal of Applied Polymer Science, 87, 2083–2099 (2003). 56. B. Baghaei, S. H. Jafari, H. A. Khonakdar, I. Rezaeian, L. As’habi and S. Ahmadian, Polym. Bull., 62, 255–270 (2009). 57. B. K. Kim, Y. S. Oh, Y. M. Lee, L. K. Yoon and S. Lee, Polymer, 41, 385–390 (2000). 58. J. Karger-Kocsis and V.N. Kuleznev, Polymer, 23, 699–705 (1982). 59. R. J. Gaymans and J. W. van der Werff, Polymer, 35, 3658–3664 (1994). 60. E. Biber, G. G¨und¨uz, B. Mavis and U. Colak, Materials Chemistry and Physics, 122, 93–101 (2010). 61. C. B. Bucknall, Toughened Plastics, Applied Science, London (1977). 62. S. Wu, J. Appl. Polym.Sci., 26, 1855–1863 (1985). 63. C. B. Bucknall and D. R. Paul, Polymer, 50, 5539–5548 (2009). 64. R. J. M. Borggreve and R. J. Gaymans, Polymer, 30, 63–70 (1989). 65. A.Yu. Vasilenko, D. D. Novikov and E. V. Prut, Doklady Physical Chemistry, 428, 2 (2009). 66. A. Gonzalez-Montiel, H. Keskkula and D. R. Paul, Polymer, 36, 24, 4587–4603 (1995). 67. G. Tripathi and D. Srivastava, Materials Science and Engineering A 443, 262–269 (2007).

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68. R. Thomas, J. Abraham, P. Thomas, and S. Thomas, Journal of Polymer Science: Part B: Polymer Physics, 42, 2531–2544 (2004). 69. S. Newman and S. Strella, J. Appl. Polym. Sci., 9, 2297 (1965). 70. T. Kietzke, D. Nehe and K. Landfester, Novel approaches to polymer blends based on polymer nanoprticles, Nature Mater., 2, 48–412 (2003). 71. S. Ritzenthaler, F. Court, E. Girard-Reydet, L. Leibler and J. P. Pascault, Macromolecules, 36, 118–126 (2003). 72. Y.S . Thio, J. Wu and F. S. Bates, Macromolecules, 39, 21, 7187–7189 (2006). 73. L. H. Sperling, Thermodynamics and kinetics of phase separation, in Interpenetrating Polymer Networks, D. Klemper, L. H. Sperling and L. A. Utracki (Eds), American Chemical Society, Washington, 77–123 (1994). 74. L. A. Utracki, Thermodynamics and kinetics of phase separation, in Interpenetrating Polymer Networks, D. Klemper, L. H. Sperling and L. A. Utracki (Eds), American Chemical Society, Washington, 77–123 (1994). 75. J. Karger Kocsis, O. Gryshchuk. and S. Schimitt, J. Mater. Sci., 38, 413–420 (2003). 76. D. Sopphiea, D. Klempner, V. Sendijarevic, B. Sutha, and K. C. Frisch, Interpenetrating polymer networks as energy-absorbing materials, in Interpenetrating Polymer Networks, D. Klemper, L. H. Sperling and L.A. Utracki (Eds), American Chemical Society, Washington, 39–75 (1994). 77. D. Klempner and D. Sophiea, Interpenetrating polymer networks, in Elastomer Technology Handbooks, N. P. Cheremisinoff (Ed.), CRC Press, Boca Raton, FL, 421–444 (1993). 78. J. Karger-Kocsis, O. Gryshchuk and N. Jost., J. Appl. Polym. Sci., 88, 2124–2131 (2003). 79. S. C. Kim, D. Klempner, K. C. Frisch and H. L. Frisch, Macromolecules, 10, 6 (1977). 80. V. Huelck, D. A. Thomas and L. H. Sperling, Macromolecules, 5, 4 (1972). 81. J. John, R. Suriyakala, S. Thomas, J. M. Mendez, A. Pius and S. Thomas, Journal of Materials Science (2010). 82. J. John, D. Klepac, M. Didovi, C. J. Sandesh, Y. Liu, K. V. S. N. Raju, A. Pius, S. Valic and S. Thomas, Polymer, 51, 2390–2402 (2010). 83. P. Bauer, J. Henning, and G. Schreyer, Angew. Makromol, Chem., 11, 145 (1970). 84. D. De Rossi, K. Kajiwara, Y. Osada. and A. Yamauchi (Eds), Polymer Gels, Plenum: New York (1991). 85. Y. Hirokawa and Y. Tanaka, J. Chern. Phys., 89, 6379 (1988). 86. E. S. Matuo and T. J. Tanaka, Chern. Phys., 89, 1695 (1988). 87. A. Suzuki and T. Tanaka, Nature (London), 346, 345 (1990). 88. Y. H. Bae, T. Okano and S. W. Kim, J. Polym. Sci., Part B: Polym. Phys., 28, 923–936 (1990). 89. K. Otake, H. Inomata, M. Konno and S. Saito, Macromolecules, 23, 283 (1990). 90. T. Okano, Y. H. Bae, H. Jacobs and S. W. Kim, J. Controlled Release, II, 255–265 (1990). 91. H. Inomata, S. Goto, K. Otake, and S. Saito, Langmuir, 8, 687–690 (1992). 92. M. Shibayama, M. Morimoto and S. Nomura, Macromolecules, 27, 5060–5066 (1994). 93. E. Kato, J. Chem. Phys., 106, 3792–3797 (1997). 94. S. Hirotsu and A. Onuki, J. Phys. Soc. Jpn, 58, 1508–1511 (1989). 95. F. Afroze, E. Niles and H. Berghamns, J.Mol.Struct., 554, 54 (2000). 96. D. Dhara and P. R.Chaterji, Polymer, 41, 6133 (2000). 97. M. Annkaka, K. Motokawa, S. Sasaki, T. Nakahira, H. Kawasaki, H. Maeda, Y. Amo and Y. Tominaga, J.Chem. Phys., 113, 5980 (2000). 98. Y. Li and T. Tanaka, J.Chem. Phys. 41, 6133 (1980). 99. L. Liang, P. C. Roeke, J. Liu, G. E. Fryxell, J. S. Young, M. H. Engelhard and K. L. Alford, Langmuir, 16, 8016 (2000). 100. T. Okajima, I. Harada, K. Nishio and S.Hirotsu, Jpn. J. Appl. Phy., 39, 1875 (2000). 101. C. Nakamoto, T. Motonaga and M. Shibayama, Macromolecules, 34, 911 (2001). 102. M. Shibayama, S. Tanaka and T. Norisuye, Physics A, 249, 245 (1998). 103. R. Yoshida, K. Uchida, Y. Kaneko, K. Sakai, A. Kikuchi, Y. Sakuraia and T.Okano, Nature (London), 374, 249 (1995). 104. Y. Kaneko, S. Nakamura, K. Sakai, A. Kikuchi, T. Aoyagi, Y. Sakkurai and T. Okano, J. Biomater. Sci. Polym. Ed., 10, 1079 (1999). 105. S. Hirotsu and A. Onuk, J. Phy. Soc. Jpn., 58, 1508 (1988). 106. T. Takigawa, H. Araki, K. Takahashi and T. Masuda, J. Chem. Phys., 113, 7640 (2000).

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Handbook of Multiphase Polymer Systems K. Haraguchi and T. Takehisa, Adv. Mater., 14, 1120 (2002). K. Haraguchi, R. Farnworth, A. Ohbayashi and T. Takehisa, Macromolecules, 36, 15, 10162 (2003). K. Haraguchi, T. Takehisa, S. Fan, Macromolecules, 35, 10162–10171 (2002). C. S. Kim and S. M. Oh, Electrochimica Acta, 46, 1323–1331 (2001). Y. Youssef and J. Denault, Polymer Composites, 19, 3 (1988). L. U. Devi, S. S. Bhagawan and S. Thomas, J Appl Polym Sci., 64, 1739 ( 1997). K. Okubo,T. Fujii and Y. Yamamoto, Composites: Part A, 35, 377–383 (2004). K. Joseph, S. Thomas and C. Pavithran, Polymer, 37, 5139 (1996). W. C. J. Zuiderduin, C. Westzaan, J. Hu´etink and R. J. Gaymans, Polymer, 44, 261–275 (2003). H. Tang, X. Chen and Y. Luo, Eur. Polym. J., 32, 8, 963–966 (1996). G. Kickelbick, Prog. Polym. Sci., 28, 83 (2003). M. Yoshida, M. Lal, N. D. Kumar and P. N. Prasad, J. Mater. Sci., 32, 4047 (1997). T. J. Pinnavaia and G. W. Beall, Polymer-Clay Nanocomposites, John Wiley & Sons Ltd., Chichester, UK, (2001). K. Friedrich, S. Fakirov and Z. Zhang, (Eds), Polymer Composites – From Nano- to Macro Scale, Springer, NY, USA (2005). G. Schmidt and M. M. Malwitz, Current Opinion in Colloid and Interface Science, 8, 103 (2003). W. Wang, Y. Li, L. Wei and Y. Fang, Mater. Lett., 57, 3366 (2003). S. Varghese, K. G. Gatose, A. A. Apostolov and J. Karger-Kocsis, J. Appl. Polym. Sci., 92, 543 (2004). S. Varghese and J. Karger-Kocsis, Polymer, 44, 4921 (2003). J. H. Lee, D. Jung, C. E. Hong, K. Y. Rhee and S. G. Advani, Compos. Sci. Technol., 65, 1996 (2005). Y. C. Ke and P. Stroeve, Polymer-Layered Silicate and Silica Nanocomposites, Elsevier (2005). W. Liu, S. V. Hoa and M. Pugh, Composites Science and Technology, 65, 307–316 (2005). Y. X. Zhou, P. X. Wu, Z-Y. Cheng, J. Ingram and S. Jeelani, eXPRESS Polymer Letters 2, 1, 40–48 (2008). A. Lee and J. D. Lichtenhan, Journal of Applied Polymer Science, 73, 1993–2001 (1999). J. T. Kim, T. S. Oh, and D. H. Lee, Polym. Int., 52, 1058 (2003). T. P. Mohan, M. R. Kumar and R. Velmurugan, J Mater Sci., 41, 2929–2937 (2006). L. F. Cai, Y. L. Mai, M. Z. Rong, W. H. Ruan and M. Q. Zhang, eXPRESS Polymer Letters, 1, 2–7 (2007). J. C. Halpin, J. Compos. Mater. 3, 732 (1969). J. C. Halpin and J. L. Kardos, Polym. Eng. Sci., 16, 344 (1976). T. Mori, K. Tanaka, Acta Metall. Mater., 21, 571 (1973). Y. Benveniste, Mech. Mater., 6, 147 (1987). T. Mura, Micromechanics of Defects in Solids, 2nd ed., Martinus Nijhoff, Boston (1987). G. P. Tandon and G. J. Weng, Polym. Compos., 5, 327 (1984). D.A. Brune and J. Bicerano, Polymer, 43, 369 (2002). T. D. Fornes, D. R. Paul, Polymer, 44, 4993 (2003).

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8 Rheology and Viscoelasticity of Multiphase Polymer Systems: Blends and Block Copolymers Jean-Charles Majest´e Laboratoire de Rh´eologie des Mati`eres Plastiques CNRS, St Etienne, France

Antonio Santamar´ıa Polymer Science and Technology Department, Faculty of Chemistry, University of the Basque Country, San Sebasti´an, Spain

8.1 Introduction This chapter contemplates fundamental results and recent advances in the rheological and viscoelastic properties of polymer blends and block copolymers that compose multiphase systems. For thermodynamic and entropic reasons, most of the polymers are immiscible and the final blend is generally a multiphase system. The microphase separation observed in block copolymers, which leads to nanoscale-ordered (self-assembled) morphologies, also lies on thermodynamic basis. From a rheological point of view, these systems require much care; rheology relations are grounded on the basic principle of continuity, homogeneity and isotropy and these principles are rarely obeyed in multiphase systems. This is the origin of all the peculiar behaviors encountered when studying the viscoelastic properties of immiscible polymer blends and microphase-separated copolymers. Homogeneity can be recovered when the length-scale of the flow applied to the sample is sufficiently larger than the size of the flow element. This allows the multiphase system to be considered as a homogeneous system, one having an average ‘specific’rheological behavior. Among the numerous influences (concentration, flow geometry, time scale, type of flow field, thermodynamic interactions between the phases. . .), morphology is the most characteristic and important parameter for such multiphase systems. The term ‘morphology’refers to the overall physical shape or physical structure of a material. For multiphase systems, such as immiscible polymer blends and microphase-separated Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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Figure 8.1 Microphase-separated block copolymers. The midblock chains (lines) constitute the elastomeric matrix. The circles indicate glass domains.

block copolymers, morphology concerns mostly the dispersed phases (spherical, ellipsoidal, cylindrical). But in addition, co-continuous phases, lamellae, fibrils, ribbons, sub-inclusions must be considered, as well as size distribution and orientation of the phases.. Actually, according to the phase dispersion state, the rheological response of the system can be radically different. Implicitly, the behavior of the system will be strongly linked to the interfacial tension between the components. The interface or interphase plays a predominant role in viscoelastic properties, particularly in the molten state of polymer blends. The total amount of interface coupled with the thermodynamic interaction between the phases is then a specific and important parameter in order to understand and control the viscoelastic properties of multiphase systems. Moreover, the current commercial interest in blends is to a large extent concerned with the dispersion and mixing of polymers in melt compounding processing equipment. A wide range of sizes and shapes can be obtained during processing. In this case, the final morphology is a balance between deformation-disintegration phenomena and coalescence essentially governed by processing conditions and materials parameters. The thermodynamic basis of binary polymer blends and diblock copolymers exhibiting phase separation have been examined in the sound morphological study performed by Bates [1]. Microphase-separated block copolymers constitute fascinating materials from both a scientific and a technological point of view. The socalled thermoplastic elastomers (TPEs), with the processability of thermoplastics and the physical properties of elastomers, are actually block copolymers with a glassy or crystalline minor phase embedded within an elastomeric matrix, as shown in the scheme of Figure 8.1. The principal aim of this chapter is to show the capabilities of rheology as a performing tool to investigate multiphase polymer systems, such as immiscible blends and block copolymers, considering the liaisons shown in Figure 8.2. For a better comprehension of the basic and applied aspects contained in this framework, the chapter is divided in two parts. The first part refers to the rheology of multiphase polymer blends in general, whereas the second part is related to the case of Poly (Styrene-b-Butadiene-b-Styrene) (SBS) triblock copolymers, as an example of a self-assembled nanostructured system. After this introduction, the chapter goes on to discuss some morphological aspects and the basic information needed to understand phase formation and characterization of immiscible blends. Then, the main materials parameters or rheological parameters influencing the morphology are reviewed. The next section describes the main and fundamental physical properties and mechanisms necessary to control and understand the morphology (dispersed-phase size). The rheology of multiphase polymer blends is then reviewed. This section presents the specificity of the rheology of immiscible polymer blends and the major models which attempt to describe the viscoelastic behavior of these systems. This part emphasizes the role of morphology on the properties in the molten state and reviews the methods developed to infer morphological information from rheological data.

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Structure Order-Disorder Microphase Separation

Morphology Viscosity

Oscillatory flow Predictive models Industrial Processing

Figure 8.2 Rheological tools to investigate basic and applied aspects of phase-separated blends and block copolymers.

The next sections are devoted to microphase-separated block copolymers. These materials are actually selfassembled systems in which an order at nanoscale is developed. Taking styrenic copolymers as a reference, it is shown how rheological techniques have turned out to be a powerful tool to characterize finely the spatial organization and phase separation of block copolymers. Finally, the chapter ends with the description of the strong impact of flow on the SEBS morphology. Flow-induced morphology is investigated for capillary flow along with the flow instabilities in SEBS copolymers.

8.2 Morphology of Polymer Blends 8.2.1

Morphology Characterization

Morphology is the source of rheological behavior in the molten state as well as for final properties. Its characterization is then essential. Many techniques have been extensively developed and used (see chapters in this book) to improve the visualization of the macroscopic and microscopy organization of the segregated phases. The goal of the following section is to give a brief description of tools which generally are the base and complete a rheological study of polymer blends. SEM and TEM visualization is not mentioned because this wide subject requires special attention. 8.2.1.1

Glass Transition

From a rheological point of view, the glass transition results in a huge decrease of the strorage modulus G of many orders of magnitude when the temperature increases and the presence of a maximum on the loss modulus G and the damping factor tan δ = G  /G  . In the case of a miscible blend (polycaprolactone/ polyvinylchloride, Figure 8.3(a)) a unique transition is observed. For immiscible blends (Polycaptolactone/ polystyrene, Figure 8.3(b)), two transitions are visible at temperatures corresponding to the glass transition of each component. However, the presence of two transitions clearly separated for a blend of two polymers only accounts for a positive criterion of nonmiscibility. The existence of a unique transition occurring over a narrow range of

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Figure 8.3 Dynamic mechanical measurements of blends with polycaprolactone.(PCL). (a) PCLpolyvinylchloride miscible blend, (b) immiscible blend PCL-polystyrene.

temperature is not a priori a proof that the system is made up of one phase. Two mains reasons are invoked. The technique fails when the gap between the glass transition of the two neat polymers is inferior to 20◦ K. Moreover, the observation of a unique but wide transition reveals a limit of miscibility. Fully miscible systems exhibit a unique and narrow transition. The second reason involved in the limitation of this technique is the contribution of the interfaces or interphases when the segregated domains are too small. Near the interface, the two species are necessarily blended. This proportion of interfacial chains increases when the characteristic size of the dispersed phase decreases.

8.2.1.2

Selective Dissolution

The aim of the selective dissolution is to get the fraction of the component belonging to the percolating structure, on both sides of the composition window. Generally, samples are firstly properly shaped as disks, then weighted, and finally immersed in a large excess of selective solvent during 5 days on an oscillating table. Solvents are carefully chosen to dissolve completely one component without any influence on the other. After extraction of the selected phase, the remaining parts are taken out of the solvent, and dried under vacuum until their weight is found to be constant. The comparison between the weight of the remaining phase and its original fraction in the blend, after and before the selective extraction, leads to the amount of continuity of the component. The continuity of one phase can be defined as the fraction of polymer that belongs to a continuous phase, which can be evaluated with the following expression: % continuity of i =

weight of the sample – weight after extraction of i weight of component i after extraction

(8.1)

Some authors [2] point out the influence of the sample size on the solvent extraction experiments. This characterization is usually applied to describe the morphology of polymer blends. In several studies [3–8], after dissolution of one component in an appropriate solvent, samples are examined with microscopy techniques. When two selective solvents are available [3, 9], both sides of the co-continuous curve are measured and thus the full co-continuity domain is determined.

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0 0

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Figure 8.4 Continuity of PEO and PVdF-HFP phases as a function of PEO content for PEO/PVdF-HFP blends. Reprinted from [9]. Copyright (2004) with permission from Elsevier.

Figure 8.4 shows the continuity curves for the two components PEO and PVdF-HFP, calculated as presented above. Both curves present the same shape. Three different zones can be distinguished. At low compositions, continuity is near to zero. As the composition increases, a percolation threshold appears describing the creation of the first connections between close droplets. In the last zone, for higher compositions, the polymer fully contributes to the 3D structure, contrary to the previous zone where some droplets still coexist with the percolating structure. In this system, PEO seems to percolate around 15%, and PvdF-HFP around 20%. The domain of co-continuity goes from 40 to 55% of PEO content. This range of full co-continuity depends on the conditions of mixing. As seen in other papers in which double selective dissolution is used, a wide range of composition for co-continuity is not surprising [6]. The curves have almost the same ‘S’ shape. However, the PVdF-HFP curve presents a singularity at low volume fractions. The continuity reaches a higher level than expected. The droplets/matrix morphology does not seem to exist. This is one of the artifacts often encountered using this technique. One of the explanations might be a partial miscibility of the two polymers near this particular composition. Another possibility could be the existence of a 1D fibrillar or highly oriented structure crossing the whole sample. Strange or wrong results are also obtained when the selective solvent of one phase causes partial swelling and partial dissolution of the other phase. The continuity calculated from the weight after dissolution is wrong. Correction can be applied to take into account the percentage of polymer unfortunately dissolved. The swelling of one phase by its nonsolvent is a more insidious problem. Swelling causes the closing of the channels needed by the dissolved polymer to flow out of the sample. The result is an underestimation of the continuity percentage.

8.2.2 8.2.2.1

Effect of Rheological Parameters on Morphology Viscosity Ratio

Viscosity ratio has turned out to be one of the most critical parameters for the control of the morphology of blends. Generally speaking, if the minor component has lower viscosity than that of the main component, the

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dispersed phase will be uniformly and finely dispersed. Reciprocally, the minor component will be roughly dispersed if its viscosity is higher than that of the main component. Avgeropolous et al. [10] proved by TEM analysis that the particle size of EPDM-BR increases with the viscosity ratio (calculated from torque ratio). Karger-Kocsis et al. [11] studied blends of EPDM-polypropylene and also noticed a linear relation between the mean diameter of particles of EPDM phase and the viscosity ratio. It is generally admitted that the smallest particle size is achieved for viscosity ratio around unity. This condition must be taken into account when a new polymer blend is developed. On the other hand, Rumscheidt and Mason [12] and Grace [13] showed that under shear flow, the droplets reach stable ellipsoid shape but do not break up for viscosity ratio above 3.7. However, a large variety of studies also showed that mixing conditions of the blends in terms of viscosity ratio can be widely different. Wu [14] proved that a caoutchoutic phase dispersed in rubber–polyamide blends can break during twin-screw extrusion even for viscosity ratio greater than 4. Favis and Chalifoux [15] showed that for polycabonate-polypropylene blends, particle break-up can occur even for torque ratio (equivalent to viscosity ratio) above 13 in a internal batch mixer (essentially a shear-mixing device). The variation of droplets size with the viscosity ratio observed in twinscrew extruder closely looks like that obtained for Newtonian fluids under elongational flow. The presence of such flow is a means to explain the results but the difficulty in quantifying the relative proportion of each type of flow does not allow rejecting other hypothesis-like viscoelasticity of the components. Nevertheless, viscosity ratio has been shown to have a strong effect on blends morphology, i.e., the size of the dispersed phase increases by factor 3 to 4 when torque varies from 2 to 13 (Figure 8.5). The minimum droplet size of the minor phase is achieved for viscosity ratio around unity (0.25 for torque ratio). This minimum size agrees relatively well with the minimum observed on Newtonian systems under shear flow. When elongational flow is present, this minimum is less marked and break-up can occur over a wider range of viscosity ratio. On the whole, these results agree with the Taylor theory [16, 17] showing that the main break-up mechanism is an almost regular step-wise reduction of the droplets by hydrodynamic stress generated by the matrix. In

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7% PC

23% PP

1 7% PP

0 10–2

10–1

100

101

102

TORQUE RATIO

Figure 8.5 The influence of torque (viscosity) ratio on the number-average diameter for blend of polypropylene and polycarbonate. Reprinted from [15]. Copyright (1987) with permission from John Wiley & Sons.

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the case where capillary instabilities play an important role, the mechanism controlling the effect of viscosity ratio can be totally different. Even in the best conditions of viscosity ratio, a high value of the viscosity of the matrix leads to a significant increase of the break-up time of the thread by capillary instabilities. The consequence is generally, for Newtonian fluids, a smaller mean size of the droplets with also a narrow size distribution. This mechanism is generally observed for very low viscosity ratio. 8.2.2.2

Elasticity

The exact role of elasticity in the determination of the size and shape of the dispersed phase in polymers blends remains one of the more misunderstood aspects. A classical study in this field has been conducted by Van Oene [18] who showed that in capillary flow, there are two main modes of dispersion:

r r

stratification shaping of droplets-fibers

These morphologies are controlled by the size of the particles, the interfacial tension and the difference in viscoelastic properties between the two phases. This study [18] proved that the elastic contribution to the interfacial tension can induce encapsulation of the low elasticity phase by the high elasticity one. Van Oene proposed to integrate elasticity effects as an interfacial effect. Then the following equation has been developed: αeff = α +

d [(N2 )d − (N2 )m ] 12

(8.2)

where α eff is the efficient interfacial tension in dynamic conditions, α is the static interfacial tension, d is the droplet diameter and N 2d , N 2m are respectively the second normal stress difference for the dispersed phase and the matrix. Elasticity of both components is taken into account through a modification of the efficient interfacial tension. Using the Couette device and different fluids with the same viscosity but different normal stress levels at room temperature, Elmendorp and Maalcke [19] studied the micro-rheological behavior of polymer blends. They demonstrated that, from an experimental point of view, the data on droplets deformation and break-up agree well with the theory for Newtonian fluids. However, when using viscoelastic fluids, one can notice that normal forces generated by the droplets tend to stabilize it, as predicted by Van Oene [18]. Using an in situ visualization device, it was shown [20] that under simple shear flow, droplets of polypropylene are stretched perpendicularly to the flow direction when the matrix of PS in more elastic. This ‘widening’effect has shown to be proportional to the second normal stress difference between the two phases. 8.2.2.3

Shear and Stress

The effect of shear rate (or stress τ12 = ηm γ˙ ) indicates that the particle size is inversely proportional to the applied shear stress. The consequence is that when increasing the shear stress, one can observe a decrease of the characteristic size of the dispersed phase. For instance, for a PS-PE blend [21], it has been noticed that applying high shear stress leads to very small droplets. In this particular case, the shear stress seems to be predominant compared to the viscosity ratio (variations of the molecular weight of PS or the temperature of the blend, i.e., all the parameters involved in the variation of the viscosity ratio, have no effect on morphology). Generally speaking, all the results seems to indicate that large variations of the shear stress are needed to

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overcome the effect of viscosity ratio in the control of dispersed phase size. If shear stress is too low, the particle size is mostly determined by the interfacial tension, viscosity ratio and elasticity. However, changing the shear stress can affect considerably the shape of the dispersed phase. The formation of fibers of PE in PS [22] matrix has been observed under uniform shear in capillary flow. The formation of fibrils may be driven by a combination between viscosity ratio and shear stress. Dealing with polymer blends in the molten state, the formation of fibril occurs for a critical level of stress. This level seems to decrease when viscosity ratio decreases.

8.3

Microrheology of Droplet Deformation

Multiphased systems differ from others because all the rheological quantities are not necessarily continuous in the volume of the sample. Then, the local viscoelastic behavior, i.e., the rheology at the microscopic scale, is prominent to describe the evolution of the morphology during the process. It is commonly admitted that the morphology of a polymer blend is mostly a balance between two mechanisms: 1. The forces generated by the matrix which tend to deform the droplets and sometimes provoke drop break-up. 2. The natural tendency of nature to minimize energy so it takes the chance to minimize the interface area by coalescence event when two droplets are colliding. Each mechanism has opposite effects on the characteristic size of the dispersed phase. Hydrodynamic stresses produced by the matrix lead to smaller droplets while coalescence results in coarse-grain morphology. Generally speaking, besides the thermodynamic interactions between the phases, mixing conditions have a strong impact on both mechanisms. We present hereafter a summary of these two elementary mechanisms. 8.3.1

Breakup

Two basic mechanisms are involved in the dispersion of one fluid in another. One is the regular step-wise equilibrium of repeated break-up. The droplet splits into two smaller parts (Figure 8.6). The other is the transient mechanism of thread break-up during extension and the disintegration into smaller droplets. This

Figure 8.6 Schematic presentation of the two main dispersion mechanisms: on the left, the step-wise equilibrium mechanism of repeated break-up at Cacrit and on the right the transient mechanism of thread breakup. Reprinted from [23]. Copyright (1993) American Institute of Physics.

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last mechanism is also called ‘capillary instability’ and is often observed during transient shear or elongational conditions of after cessation of flow.

8.3.1.1

Regular Step-wise Break-up Above Critical Conditions

When studying Newtonian fluids suspended in another Newtonian fluid, Taylor [16, 17] observed that when the radius of the droplets is great enough or when the shear rate is high, the droplets break up. He proposed a series of equations from which one can obtain the size of the smallest drop existing in the fluid whatever the shear rate. The parameters governing the break-up are the viscosity ratio, the type of flow, and the so-called capillary number Ca which is the ratio between the deforming stress ηm γ˙ and the interfacial stress α/R where α is the interfacial tension and R the radius of the droplet. Taylor proposed the following equation for the capillary number Ca: Ca =

ηm γ˙ α/R

(8.3)

For small values of Ca, the interfacial forces dominate and a steady drop shape develops. Above a certain critical value, represented by Cacrit , the drop becomes unstable and finally breaks. This simple equation is the basis of the understanding of break-up mechanisms in polymer blends. Eq. (8.3) defines clearly but naively the conditions for droplets break-up. Experimental observations show more complex situations. For instance, Eq. (8.3) does not involve viscosity ratio whereas it was shown in previous sections that this parameter is predominant. Moreover, Taylor does not include coalescence mechanism. This theory can be used to predict the morphology of polymer blends undergoing pure break-up mechanism. This is of course an ideal situation. In order to reach a realistic prediction of the morphology, it is preferable to combine pure break-up and pure coalescence mechanism provided that no strong interaction exists between the two mechanisms. The use of capillary numbers has been generalized by Grace [13] who performed measurements on blends of two Newtonian polymers fluids. They studied two type of flow: uniform shear flow and hyperbolic flow (equivalent to elongational flow). Elongational flow proves its efficiency to induce break-up of the droplets. The viscosity ratio range – for which the droplets can break – seems to be much larger than for shear flow. Figure 8.7 shows the conditions of break-up as a function of the viscosity ratio. Droplets break up for Ca > Cacrit and are stable (with equilibrium shape) when the capillary number is below the critical value. On Figure 8.7, the curve corresponding to simple shear flow is an empirical fit to Grace’s data by De Bruijn [24]. The following tractable equation has been found: log Cacrit = −0.506 − 0.0995 log p + 0.124(log p)2 −

0.115 log p − log 4.08

(8.4)

Under shear flow, the deformation and breakup is easiest for viscosity ratio from 0.25 to 1. Under elongational flow, the curve is flatter, indicating that conditions for droplets breakup are less rigorous. For Newtonian droplets in a Newtonian matrix under shear, Karam and Bellinger [25] and Tavgac [26], found typical U_type dependence of droplet break-up on the viscosity ratio, which indicates that upper and lower limits exist for viscosity ratio beyond which no particle breakup can occur. The range of viscosity ratio where break-up is observed may change from one author to another [25, 26] (0,005 ≤ p ≤ 3,0, 0,0033 ≤ p ≤ 3,7.) but the shape of the curve remains the same. As a rule, for high viscosity ratio, the viscous forces that

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1000 Simple shear 100 Cacrit

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1

Elongation

0.1

0.01 1e–5

1e–4

1e–3

1e–2 1e–1 viscosity ratio

1e+0

1e+1

1e+2

Figure 8.7 Critical capillary number versus the viscosity ratio p = ηd /ηm in simple shear and elongational flow. The curve for the shear flow is Eq. (8.4). Reprinted from [13]. Copyright (1982) with permission from Taylor and Francis.

disrupt the drop cannot overcome the interfacial forces that ensure the cohesion of the drop. By contrast, at very low viscosity ratio, the droplet is very deformed without break-up.

8.3.1.2

Capillary Instabilities

Since a sphere offers the minimum surface area for a given volume, the interfacial tension between two fluids causes any non-spherical shape for a volume of fluid suspended in another to be unstable. A capillary instability is defined as the instability observed for a cylindrical thread of fluid in another fluid. Rayleigh [27] was the first to suggest a model for the growth of a disturbance in a viscous jet in air. Later, Tomotika [28] extended this theory to Newtonian fluids. He noted that the time required for thread break-up through capillary instability depends on parameters such as interfacial tension, viscosity ratio and initial thread diameter. The cylindrical thread will deform gradually and sinusoidally (Figure 8.8). According to theory, distortions with a wavelength λ greater than the initial circumference of the thread will grow exponentially with time: ε = ε0 exp(qt)

(8.5)

ε Ro

λ

Figure 8.8 Schematic presentation of capillary instability appearing during extension of a thread with intial radius R0 .

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where ε is the amplitude of the distortion at time t, ε0 is the initial distortion and the grow rate q of this distortion is given by: q = α (λ, p)/2ηm R0

(8.6)

where α is the interfacial tension, ηm is the zero shear viscosity of the matrix and R0 is the initial radius of the thread. The function can be calculated form Tomotika’s original equations and has a maximum for a given wave number defined by Xm = 2π R0 /λm . This number depends on the viscosity ratio. Then, for a given viscosity ratio, there will be a dominant wavelength for which the amplitude grows faster. Tomotika estimated that the thread break-up will occur when ε = 0.8R0 . From Eq. (8.5), the break-up time tb can be estimated using the following expression: tb = ln(0.8R0 /ε0 )/q

(8.7)

The Tomotika [28] model has successfully described experimental observations of quiescent breakup. Other authors have shown some unpredicted effects such as satellite droplets [29] droplets and ‘end picking’ [29]. The effect of elasticity has been demonstrated experimentally. It seems that instabilities initially grow faster but break-up is delayed by extension thickening at large strain [30]. It is worth noting that the measurement of capillary instabilities could be used to infer interfacial tension in immiscible polymers blends. 8.3.2

Coalescence

Coalescence is the natural tendency of multiphase system to minimize the interfacial area and thus the free enthalpy. This mechanism occurs when hydrodynamic forces are not able to provoke the rupture of the droplets, i.e., at low capillary number. Consequently, characteristic droplet size increases. During the mixing of polymers, rupture and coalescence interact and finally balance, thus giving the equilibrium morphology that will govern the final mechanical properties. As this morphology is a balance between two mechanisms, changing the mixing condition will obviously change the final state. This phenomenon called ‘morphological hysteresis’ have been demonstrated experimentally by Minale et al. [31, 32]. Many authors (Chesters [33], Janssen [30], Ottino et al. [34]) have reviewed the coalescence model. From all these works, it appears that the critical event in coalescence is the collision between droplets. Without collision there is obviously no coalescence but a collision of two drops will not always end in a unique bigger drop. Colliding droplets develop a flat interface separated by a thin film. On one side hydrodynamic forces push the drop together and on the other side, the thin film can resist or rupture. Coalescence will not occur if the film resists or if the flow reverses and separates the two droplets. Thus it appears that coalescence has to be considered as a dynamic phenomenon [35]. The contact time (or interaction time) between the droplet is a crucial parameter. Consequently, coalescence mechanism will be promoted by low shear rate and low resistance films between the drops. The resistance of the film is governed by the time needed by the matrix to be drained between the contacting droplets. The literature mentions two limit approximations for film drainage. One model considers a fully mobile interface. It seems to be the case for blends with very small viscosity ratio. The other model assumes immobile interface which is appropriate when the viscosity ratio is largely greater than unity. Each drainage model gives a different dependence of the coalescence-limited drop size on shear rate. Partial mobile interface has been also introduced for intermediate viscosity ratio. Both models describe the evolution of the droplet size with time. Frequency of collision, interaction time and hydrodynamic forces are obtained through scaling arguments and finally are combined with the film drainage model to give the probability that a collision results in a coalescence event. The time evolution

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of droplet size is governed by a differential equation involving parameters like the shear rate, the volume fraction, the initial drop radius, the interfacial tension. Both models predict the time evolution of the drop size until a pseudo-steady drop size Rss is reached. The size Rss is associated with the current mixing conditions (shear rate, shear stress, capillary number). For the immobile interface theory, the pseudo-steady radius Rss scales as γ˙ −1/2 . For the fully mobile interface, authors (Janssen [30] and Minale et al. [31, 32]) agree to find Rss ln Rss ∝ γ˙ −1 . The interesting point is the use of such scaling law to understand the correlation between the final morphology and the mixing conditions. The pseudo-steady drop size under pure break-up mechanism scales as γ˙ −1 . Tucker and Moldenaers [35] demonstrated the morphological hysteresis plotting two curves in the space of the drop diameter versus the shear rate. One for the limit of the coalescence mechanism (immobile interface Rss ∝ γ˙ −1/2 ), the other for the break-up limit (Rss ∝ γ˙ −1 ). The space is then separated into many regions. Thus they showed that depending on the values of the couple {shear rate, drop size} the morphology is stable or presents the morphological hysteresis phenomenon. They also pointed out the existence of a critical shear rate above which the steady-state drop size is determined by a competition between break-up and coalescence and is only a function of shear rate. Below this critical shear rate, there exists a range of droplets diameters that are too small to break up but also too large to coalesce during the mixing time. This leads to morphological hysteresis. In summary, coalescence must be considered as a dynamic mechanism in which rheological conditions play a large part.

8.4

Rheology of Polymer Blends

8.4.1 8.4.1.1

Specificity of Blend Rheology Non-uniformity of Strain and Shear-induced Segregation

In polymer blends, as in filled polymers, a heterogeneous distribution of the strain (or strain rate) in the measurement device can cause segregation of the less viscous polymer in the high strain (or strain rate) zones in order to minimize the pressure drop or the dissipated energy. Figure 8.9 shows the difference of apparent viscosity between measurements performed using cone and plate device (uniform shear) or capillary rheometer (maximum shear at the wall). Measurements done with a cone and plate geometry keep the initial state of dispersion in the system. The difference vanishes at high shear rate because the viscosity gap between the two polymers fades away. This phenomenon is illustrated for PA66/PET blend in Figure 8.10. TEM observation of the cross section of an extrudate shows the difference of concentration between the center and the edge when the shear is low. PET droplets have migrated towards the high shear rate zone in order to reduce the shear stress. Usually, this effect is more pronounced for filled polymers because of higher viscosity ratio. In addition, even for homogenous flows, multiphase systems may exhibit strong strain heterogeneities. This is mainly caused by differences in rheological behavior between the phases. For instance, when dispersed phases are more viscous than the matrix, the particles are less deformed (or totally undeformed). The macroscopic strain is then concentrated into the matrix. The local strain (microscopic strain) between the drops is greater than the strain applied to the whole system. One of the consequences is an apparent shift of the shear thinning behavior towards the low shear rate. This problem is particularly important in the case of filled semi-crystalline polymers. The onset of the stress-induced crystallization process is different compared to that obtained in quiescent conditions. Besides, the classical display of rheological curves versus the strain or strain rate must be used with care. Indeed, while the strain spatial distribution is not homogeneous, the stress remains the same in the whole volume. So, it is more convenient to plot the variation of rheological function versus the stress. In the same manner, it is preferable to plot G = f (G ) rather than G (ω) and G (ω). Figure 8.11 illustrates the differences

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104 280 °C

102 s–1

Viscosity [Pa.s]

300 °C 103 280 °C

104 s–1

300 °C

280 °C

102

105 s–1

300 °C

101

0

20

40

60

80

100

wt-%

Figure 8.9 Viscosity of PA66-PET blends versus composition, temperature and shear rate. (o: cone and plate rheometer; •: capillary rheometer).

observed for two different plots of PS/PMMA blends. Only the first normal difference variation are sensitive to the change of representation. This is due to the strong contribution of interfacial tension to the overall stress and elasticity; besides, viscosity is not so affected by morphology and the amount of interfacial area. 8.4.1.2

Morphology Flow Modifications

The study of steady shear viscosity of polyoxymethylene/copolyamide blends [36] (Figure 8.12) highlights some anomalies in the change of viscosity for high shear stress. These anomalies are linked to the

Figure 8.10 TEM microphotographs of the edge of an extrudate of PA66_PET blend after low shear rate extrusion. It clearly shows a non-uniform distribution of PET droplets in the vicinity of the edge of the extrudate.

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η[Pas]

N1[Pa]

104

104

105 103

102 0

50

W[%] 103 0 100 (a)

105

103

104

50

W[%] 100

N1[Pa]

104

102 0

W[%] 103 50 100 0 (b)

W[%] 50

100

Figure 8.11 Viscosity and first normal difference of PS/PMMA blends at 200◦ C. (a) at different shear rate, (b) at different shear stress. Reprinted from [18]. Copyright (1972) with permission from Elsevier.

transformation of the dispersed phase into fibrils. Fibrillation decreases inter-particles’ interactions and consequently the viscosity. Generally, morphologies could be very different for low shear measurements compared to relative high shear measurements. These morphological changes induce deviation from the Cox-Merz rule. Figure 8.13 shows that under steady shear flow, the viscosity deviates from the Cox-Merz rule for the higher shear rate which is consistent with Taylor theory. Besides, dynamic measurements performed on pre-sheared sample do not match the Cox-Merz rule too. Morphologies are different between the two tests. By contrast, the agreement with capillary rheometry is obtained. This effect is due to a non-reversible morphological change during pre-shear.

η[Pa.s] 2400

2000

1600

1200

800

400

0

0

20 40 60 80 100 W% CPA

Figure 8.12 Shear viscosity at 190◦ C for blends of polyoxymethylene (POM) and copolyamide/(CPA) for various wall shear stress. (104 N/m2 ): (•) 1.27; () 3.93, () 5.44, () 6.30, (o) 126, () 193, ( ) 316. Reprinted from [36]. Copyright (1975) with permission from John Wiley & Sons.

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107 (c)

Dynamic Steady shear (cone and plate) Dynamic after pre-shear Steady shear (capillary rheometer)

η, η ∗ [Pa.s]

106

105

104

103

102 10–3

Figure 8.13

8.4.1.3

10–2

10–1

100

101 . ω,γ [s–1]

102

103

104

105

Morphological modification induced by pre-shear and effect on the Cox-Merz rule.

Swelling

Generally speaking, important synergistic effects are observed due to the shape recovery of fibrillated droplets (Figure 8.14). However, for dispersed phases more rigid than the matrix (high viscosity ratio), swelling can decrease because we catch with the case of filled polymers (non-deformable droplets). It is evident that the intensity of the swelling is linked to the ability of the droplets to be deformed and above all their ability to recover from this deformation. The more deformed the particles are deformed, the more important the swelling. Once again, interfacial effects are directly involved in such phenomenon. As seen in the previous section, under flow, a droplet breaks up until equilibrium is reached. If the shape relaxation

1.5

Swelling

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1.3

0

20

40

60

80

100

wt-%

Figure 8.14

Swelling of polymer blends: ( ) LLDPE/LLDPE, (O) LLPDE/LDPE.

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time of droplets is long, they could remain in a more or less deformed state and store elastic energy. This energy is proportional to interfacial area. After cessation of flow or strain/stress release, the equilibrium is disrupted and the drops will recover their spherical stable shape without external stress. The shape recovery of the dispersed phase leads to swelling in the transverse direction.

8.4.2 8.4.2.1

Blending Laws and Viscoelasticity Models Blends of Miscible Liquids

Miscible polymer blends are less common than immiscible ones. The miscibility is generally confined to the molecular scale. Nevertheless, the literature [37] shows that numerous polymeric mixtures were identified as miscible. However, the rheological studies of miscible blends are relatively rare. In miscible blends, the flow behavior depends on free volume, entanglements and specific interactions. Assuming free volume additivity, some basic semi-phenomenological laws have been proposed for the zero shear viscosity of miscible polymers blends log(η0 ) =



χi log(ηi )

(8.8)

i

log(V η0 ) =



χi log(Vi ηi )

Glasstone et al. [38]

(8.9)

i

χ 1 = η0 ηi i  log(η0 /ρ) = χi log(ηi /ρi )

Bingham [39]

(8.10) (8.11)

i

log(η0 ) =



Wi log(ηi )

(8.12)

i

where χ is the molar fraction; W is the weight fraction; η is the zero shear viscosity; V is the molar volume; and ρ is the specific mass. Polymer blends without hydrodynamic interaction generally follow laws given by Eqs. (8.8) to (8.12). For samples without phase segregation, a semi-empirical law [40] is found under the following form: η0α



Wi ηiα

with 0.01 < α < 0.57

(8.13)

i

This previous relation predicts a positive deviation compared to the additive mixing rule. The development of the theory of polymer dynamics, and particularly the reptation model [41], provides simple mixing rules for miscible systems which prove useful to describe such systems. Generally, the model and theories deal with polymer species of the same nature; therefore thermodynamic interactions are neglected. Many models [42–44] have been proposed that take into account microstructure, distance to the glass transition temperature, topology of the molecules and many other molecular parameters. However, very few systems have been compared to these models. Some authors showed that Time-Temperature Superposition is verified [45] for such miscible blends even when specific interactions (H bonding) are present [46].

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327

Suspensions of Rigid Particles in Viscoelastic Media

Suspensions are systems for which the deformation of the dispersed phase can be neglected. When droplets are in the solid state, blends of immiscible polymers can be considered as suspensions. However, even if the rheological properties are similar, solid polymeric droplet cannot be totally considered as fillers. The multiplicity of interfaces and structure of inorganic filler considerably widen the panel of rheological behaviors. Depending on the nature of the surface, the size, the concentration or the interaction with the matrix, filled polymers exhibit particular behaviors which are not present for polymers blends with rigid dispersed phases. Filled polymers often exhibit yield stress when the percolation threshold is overcome. Particles–particles interactions are essentially involved in such specific behavior. The overall rheological behavior can be considered as the sum of two contributions: Hydrodynamic interaction between particles and particles–particles interactions. For polymers blends, we only consider the first contribution. In the following paragraphs, we will describe the models able to predict or just describe the viscoelastic behavior of suspensions. Different regimes of concentration will be described, from diluted to concentrated suspensions. 8.4.2.3

Dilute Suspensions

The collection of equations described in the literature is built on the same base: the historical work of Einstein [47] on dilute suspension of monodisperse spheres in Newtonian fluids. Einstein proposes the following equation for the shear viscosity variations of a filled fluid: η0 (φ) = η0 (0)(1 + [η]φ)

(8.14)

where [η] is the intrinsic viscosity of the modisperse spheres. Einstein found [η] = 2.5. Equation (8.14) is based on the additional viscous energy loss due to the move of nondeformable spheres in the liquid. It assumes no particles anisometry and no interparticular hydrodynamic interactions. The relation is then valid for very dilute suspensions. For practical purposes, deviation from the Einstein law indicates the onset of a semi-diluted regime (hydrodynamic interactions without particle–particle interactions). For anisometric particles, the value of the intrinsic viscosity [η] is greater than 2.5: up to 12 for prolate ellipsoids and up to 8 for oblate ones [48]. 8.4.2.4

Concentrated Suspensions

As the dispersed phase concentration increases, the velocity field perturbation around a spherical particle may interact with that of another particle. The effects of hydrodynamic interparticular interactions can be taken into account using phenomenological additional terms of degree more than unity in the Einstein relation.  η0 (φ) = η0 (0) 1 +



 ai φ

i

(8.15)

i

Terms ai (i > 2) allow the description of special effects like the size and the interactions between particles. The coefficient in the polynomial expression can change depending on the anisometry of the particles. The number of parameters in the series may be a problem, so Eq. (8.15) is often reduced to [49]: η0 (φ) = η0 (0)(1 + βφ)μ with various values for the parameters β and μ according to the authors (Table 8.1).

(8.16)

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β

Einstein Hess-Vand De Bruyn

2.5 –2.5 –1.73

μ 1 –1 –2

As shown in Table 8.1, values of the coefficients are very different according to the authors. This reflects some lack in the generalization of Einstein’s equation using polynomial development. Actually, generalization for higher concentrations quickly becomes very complicated. Indeed, hydrodynamic interactions increase considerably and it is a tricky task to determine precisely the expression of the disturbed velocity field near the particles. Generalization is therefore empirical or semi-empirical. The different models share two characteristics:

r r

The limit of models for low concentration is the Einstein relation In agreement with experimental observations, viscosity must diverge for φ = φ m the maximum packing fraction.

Among all the empirical expressions, the most used is still the model of Krieger and Dougherty [51]. The model is based on physical arguments and never introduces specific hypotheses about the granulometry. The main idea is that each particle evolves in an effective environment (composed by the suspending liquid and the other particles) which behaves as a homogeneous fluid whose viscosity only depends on concentration. In other word, local energy dissipation must match the macroscopic one. Krieger and Dougherty showed that the effective viscosity can be cast in the following general form:   φ −z ηeff (φ) = η(0) 1 − A

(8.17)

Theory imposes A = 1 but the formula has been generalized in order to introduce the maximum packing fraction. The coefficient z is fitted to experimental data and generally close to 2. To recover the Einstein formula, the coefficient A must be equal to φ m and z = 2.5φ m . Many authors [49, 50] proposed the use of the notion of particles intrinsic viscosity under the following form:  [η] = lim

ϕ→0

η0 (φ) −1 η0 (0)

 (8.18)

Thus, Krieger and Dougherty [51] have written relation 8.17 under the well-known form:   φ [η]φm η0 (φ) = η0 (0) 1 − φm

(8.19)

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Implicitly, this law assumes that particles are homogeneously distributed in the domain. This may not be the case for suspensions in low viscosity fluids for which particles segregation is often observed during experiments. Relation (8.19) is often used to infer some structural information on particles in filled fluids through the intrinsic viscosity and the packing fraction. Generally, maximum packing fraction varies from 0.52 for monodisperse sphere up to 0.8 for bimodal population of spheres. Anisometry has a strong effect on φ m so that very different ranges of values can be found for particles with a high aspect ratio. However, particles–particles interactions are not taken into account in the Krieger-Dougherty expression and this kind of interaction makes Eq. (8.19) unusable without previously subtracting the particle network contribution. 8.4.2.5

Uniform Description of Flow Curves

Viscoelastic behavior of the suspending fluid is usually taken into account through the use of reduced viscoelastic function. Kitano and Kataoka [52, 53] verified that Eq. (8.19) is satisfied in the case of polymers using relative function for polystyrene and polyethylene filled with silica beads. According to the majority of authors, the use of reduced parameters must be performed at constant stress and not at constant strain rate. Consequently, if Eq. (8.19) holds, without particle–particle interactions, the curves of one viscoelastic variable at different filler concentrations must superimpose with the curve of the neat matrix. Agoudjil et al. showed [54] that dynamic viscosity η* (ω) of EVA filled with glass beads (Figure 8.15a) follows the general equation: η(ω, φ) = η(ω, 0) f (φ)

(8.20)

where f (φ) is a general function (Einstein, Krieger-Dougherty. . .). The single curve obtained is presented in Figure 8.15b. For concentrations below 31%, the overlay of the curve is complete for the overall frequency range. For filler content higher than 31%, a significant divergence of the complex shear viscosity is observed for the lowest frequencies. This behavior is due to the presence of filler–filler interactions, a consequence of a three-dimensional structure inside the composite [55]. These types of interactions do not scale in the same way as the hydrodynamic ones. Moreover, it affects the dynamic viscosity only at low frequencies while hydrodynamic interactions are present over the whole range of frequency. They also show that shifting the curves and overlaying them with the matrix behavior allows both estimating f (φ) and finding in which frequency range Eq. (8.20) is verified. Using this method enables to check when the filler–filler interactions dominate standard hydrodynamic interactions which are the source of the increase of the viscosity. Typical experimental data of f (φ), which actually represents the variations of the composites viscosity normalized to the matrix viscosity or in other words the relative viscosity, are shown in Figure 8.16. More generally, some authors [49, 56] suggest that an examination of the flow behavior of hard or soft spheres in polymer or simple solvents can be used to deduce the role of the filler structure on the rheology of dispersions in polymer melts. They consider that the semi-empirical Krieger–Dougherty relation (Eq. (8.19)) can be used to predict how the suspension viscosity of hard spheres asymptotes towards the liquid–solid transition at maximum packing (φ m ). Important information on the filler structure can be inferred from such measurements. Recently, with the coming of new grafted fillers, this approach allows determining parameters such as the effective packing fraction, the mean distance between particles and then, for instance, the brush

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105 η ∗ (Pa.s)

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21.4% 3.6% 0%

103 10–2

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41.4%

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102 10–2

10–1

100 ω (rad.s–1)

101

102

Figure 8.15 Experimental variations of complex shear viscosity η of EVA filled with silver coated glass particles. (a) Data obtained at various fillers load; (b) single curve obtained according to Eq. (8.20). The particles network signature is visible at low frequency for the highest compositions. Reprinted from [54]. Copyright (2008) with permission from Elsevier.

thickness on the particles. Some authors [56] also try to extend the use of this relation to the case of soft particles (Figure 8.16).

8.4.2.6

Emulsions of Newtonian Liquids

Emulsions of Newtonians fluids exhibit a pronounced elastic contribution at low shear (or shear rate).This is the case for mayonnaise or some French dressings, which exhibit strong elastic behavior after mixing. This unexpected elasticity originates from interfacial forces. These forces tend to minimize the surface between the

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104

ηrel,0

103

102

16nm-2k 16nm-13k 100nm-2k 100nm-13k 600nm-8k Krieger-Dougherty

Hard sphere behavior

101 Soft sphere behavior 100 0.0

0.2

0.4

0.6 φeff

0.8

1.0

1.2

Figure 8.16 Relative viscosity versus effective volume fraction. Data show the increasing divergence of hard spheres behavior. The solid line is the Krieger-Dougherty fit with φ m = 0.63 and the dashed line is to guide the eyes. Reprinted from [56]. Copyright (2009) with permission from the Royal Society of Chemistry.

two liquids and in the same way they balance the droplets deformation during steady state flow. Consequently, an emulsion of two Newtonian liquids can exhibit a viscoelastic behavior. Taylor [16, 17] was the first to propose an analysis applicable to emulsions of monodisperse spheres of Newtonian liquids. The constitutive equation in linear viscoelasticity can be cast into the following form: τ = η M γ˙ −

Rφη2M (19k + 16)2 d γ˙ 80α (k + 1)2 dt

(8.21)

Under dynamic solicitations, the previous equation becomes: Rφηm2 (19p + 16)2 2 ω 80α ( p + 1)2   5p + 2 φ ω G  (ω) = ηm 1 + 2p + 2 G  (ω) =

(8.22a) (8.22b)

where p is the viscosity ratio; R the radius of droplets; α the interfacial tension; φ the volumic fraction of droplets. When p tends to infinity (rigid spheres), the Einstein equation is recovered. This model predicts that the zero-shear viscosity of the emulsion increases whatever the nature of the dispersed phase. At low frequency a relaxation is also predicted. However, Eq. (8.22) indicates that the storage modulus increases monotonically until infinity. Experimental observations show that the elastic effect levels off at high frequencies, which is not captured by the Taylor expression. By also considering this kind of droplets relaxation, Oldroyd [57] proposed Eq. (8.23). The behavior is then described by three constants: the zero-shear viscosity of the emulsion η0 , a relaxation time λ1 and a

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retardation time λ2 ; τ + λ1

dτ d γ˙ = η0 γ˙ + η0 λ2 dt dt

(8.23)

Under dynamic solicitations, Eq. (8.23) gives the real and imaginary part of the complex shear modulus: η0 (λ1 − λ2 ) 2 ω 1 + (ωλ1 )2 η0 (λ1 λ2 ω2 + 1) ω G  (ω) = 1 + (ωλ1 )2 G  (ω) =

The parameters cited above are:   5p + 2 5(5 p + 2)3 2 φ+ η0 = η m 1 + φ 2p + 2 10( p + 1)2   (19 p + 16) R (19 p + 16)(2 p + 3) 1+ φ λ1 = ηm α 40( p + 1) 5( p + 1)(2 p + 3)   3(19k + 16) ηm R (19k + 16)(2k + 3) 1− λ2 = φ α 40(k + 1) 10(k + 1)(2k + 3

(8.24a) (8.24b)

(8.25a) (8.25b) (8.25c)

Actually, Oldroyd’s equation had been initially formulated into nonlinear terms with a more general form for the derivative. The general equation predicts both pseudoplastic behavior and normal shear stresses. Under dynamic solicitation, the elastic modulus G exhibits a plateau at high frequencies which can be used as an indicator of the strength of interfacial forces. 8.4.2.7

Dilute and Concentrated Emulsions of Viscoelastic Polymers

The transition to a viscoelastic component is not so trivial because of the elastic nature of polymers which interacts with the interface elastic contribution. Much phenomenological modeling has been proposed to describe the behavior of the blend from the viscoelastic behavior of each phase. The oldest works are without doubt those of Kerner [58] in the case of dilute suspensions of a polymer in a polymeric matrix. Both the dispersed phase and the matrix are viscoelastic media. Kerner assumes perfect adhesion and no interfacial forces between the phases. The equation proposed by Kerner is: G ∗ (ω) = G ∗m (ω)

3(1 − φ)G ∗m + (2 + 3φ)G ∗d (3 + 2φ)G ∗m + 2(1 − φ)G ∗d

(8.26)

Originally, the previous equation was developed for a dilute solution, i.e., without considering strong hydrodynamic interaction. Lewis and Nielsen [59] have improved and completed this approach by taking into account these interactions through the maximum packing fraction. Then they changed the volume fraction in an efficient fraction which considers the maximum packing volume and then the hydrodynamic interaction appearing when concentration increases. They proposed the following equation: 

φeff

1 − φm =φ 1+ φm2

 (8.27)

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106 G', G'' [Pa]

104

102 10–3

Figure 8.17

ω [rad/s] 10–1

101

103

Best fit of experimental data for PMMA/PS blend at 200◦ C with Palierne’s model (Eq. (8.28)).

Practically, the Kerner model is often used to describe or predict the viscoelastic behavior of two viscoelastic polymers when interfacial contribution can be neglected. This is the case at high frequency for dynamic shear measurements or a low strain for nonlinear flow when the droplets are not stretched enough to develop significant interfacial forces. 8.4.2.8

General Expression for Viscoelastic Emulsions and Suspensions

Palierne [60] proposed a generalized approach of the nonmiscible polymer blends (Later, Bousmina [61] proposed asimilar model, extending the Kerner Model). He developed a viscoelastic emulsion model considering the behavior of the phases, the morphology and the interfacial tension between the phases. He gave the following set of equations: 1+3 G ∗ (ω) = G ∗m

1−2

i

φi Hi (ω) φi Hi (ω)

(8.28a)

i





 α ∗ 4 2G m (ω) + 5G ∗d (ω) + G ∗d (ω) − G ∗m (ω) 16G ∗m (ω) + 19G ∗d (ω) Ri Hi (ω) = 



 α ∗ G m (ω) + G ∗d (ω) + 2G ∗d (ω) + 3G ∗m (ω) 16G ∗m (ω) + 19G ∗d (ω) 40 Ri

(8.28b)

where G ∗d , G ∗m are respectively the viscoelastic behaviours of the dispersed phase and the matrix; φ i are the volume fraction of droplets with radius Ri ; α is the interfacial tension. Graebling et al. [62] have shown that, in most cases, the droplet size distribution {Ri , φ i }, initially implemented in the model can be replaced by the volume-average drop size Rv without loss of accuracy. Then the summation over the size distribution can be suppressed in order to give a more tractable equation. Figure 8.17 gives an example of the good agreement between model prediction and experimental behavior. The generalist nature of the Palierne model can be easily checked. The model includes all the previous equations (Einstein, Oldroyd, Kerner, etc. . .):

r r r

if G ∗d = jωηd et G ∗m = jωηm , one can recognize the Oldroyd model (emulsions of two Newtonian fluids with interfacial tension) if α = 0, G ∗d = G d et G ∗m = jωηm , one can find the Einstein model for dilute rigid spheres in Newtonian fluids if the interfacial tension is set to zero, the Kerner model is recovered (viscoelastic phases with perfect adhesion).

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The ability of the model to predict consistently the viscoelastic behavior of polymer blends from the pure phase behavior gave some authors the opportunity to use the model as a characterization tool of the blend morphology. For instance, Friedrich et al. [63] have shown that the Palierne model can be used to measure the droplet size distribution. Coupling dynamic measurements with TEM observations, they demonstrated that both techniques lead to similar size distribution. Regardless, they gave some limitations to the use of the Palierne model in such applications. The size distribution must not be too broad and the concentration of the dispersed phase must stay below 20 vol-%. This is the consequence of one of the original hypothesis in the model: the morphology must be droplet into a continuous matrix. Yet depending on the system, more complex morphologies or even co-continuous morphologies can be encountered. In these cases, Palierne’s model fails. With the intention to measure the interfacial tension, Elias et al. [64], and Vermant et al. [65] also used the Palierne model to infer the value of interfacial stress α/Rv . Using SEM or TEM observations to measure the volume-average drop size Rv , they have obtained a value of an apparent or efficient interfacial tension. Elias et al. [64] have fitted experimental data of an imaginary part of the complex shear viscosity η (ω) with the Palierne model. The interfacial tension has been adjusted to match the whole curve of η (ω) and its maximum. Vermant et al. [65] measured the experimental droplet relaxation time (at the maximum of η ) and deduced the best value of α/Rv to match with the theoretical expression given by the Palierne model. Vermant et al. preferred to use the ratio α/Rv (interfacial elasticity) to avoid any ambiguity concerning the interfacial tension. However, Elias et al. [64, 66] have shown that the value of interfacial tension measured from dynamic measurement agrees well with that obtained from more direct determination (isolated droplet relaxation after step shear). Elias et al. [66] have shown that, in some cases, the droplet relaxation time is too similar to the one of the matrix. The two peaks corresponding to each relaxation process are much more convoluted and the determination of the relaxation times becomes inaccurate. In order to overcome this difficulty, they address an improved route to calculate the ratio α/Rv . This is accessible after isolating the droplet contribution to the complex relaxation modulus. The method is based on the fact that the Palierne model can be divided in two contributions according G ∗Palierne = G ∗Composition + G ∗droplet

(8.29)

where G ∗Composition is the complex shear moduli of the blend without any interfacial effects. The Kerner model is able to predict precisely the viscoelastic properties of a blend of two viscoleastic polymers without any interfacial effects. G ∗droplet captures the interfacial effects and specially the extra elasticity brought by the droplet deformability (Oldroyd model). It is easily shown that the storage part of the complex shear modulus G for a blend of two viscoelastic fluids can be expressed as follows:



G Palierne α, G ∗m , G ∗d = G  Palierne 0, G ∗m , G ∗d + G  Palierne (α, ηm , ηd ) = G Kerner + G Palierne (α, ηm , ηd )

(8.30)

where G ∗m , G ∗d , are respectively the complex shear moduli of the matrix and the dispersed phase; ηm , ηd are respectively the Newtonian viscosity of the matrix and that of the dispersed phase. In the previous equation, only the second part depends on the interfacial tension. This corresponds to the droplet contribution G ∗droplet . This part can be easily isolated. The main idea of the method is to reveal the contribution of the relaxation of the droplets which depends on the interfacial tension α by subtracting the composition effects to the experimental data of G* . This implies obviously that Kerner’s model correctly describes the polymer blend complex shear modulus. This is checked in the high frequency zone for which the

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106

105

G', G'' (Pa)

104

103

102 Kerner Model Palierne Model G' G''

101

100

10–1 –2 10

10–1

100

101

102

103

ω (rad/s)

Figure 8.18 Modeling of the viscoelastic behavior of PP/EVA03 80/20 with 3 wt% silica particles. The thick line represents the prediction from the Kerner model (Eq. (8.26)), T = 200◦ C. The thin line corresponds to the Palierne model (Eq. (8.28)) with α/R = 550 Pa. Reprinted from [66]. Copyright (2008) with permission from John Wiley & Sons.

effect of droplet relaxation and then interfacial tension is negligible (Figure 8.18). According to this method, the effective interfacial tension of PP/EVA blend was derived from the experimental measurement of Rv and fitting the experimental curves of the complex shear modulus [66] (Figure 8.18). On this figure, the droplet contribution is visible at low frequency characterized by an additional relaxation on G . Even if the high frequency zone is well described by both the Kerner and Palierne models, only the Palierne expression is able to take into account the interfacial forces and to predict the overall behavior.

8.4.3 8.4.3.1

Low Frequency Viscoelastic Behavior of Polymer Blends Interfacial Area and Morphology

Studying viscoelastic properties seems to be a suitable way to characterize and understand polymer blends. Many authors [3, 4, 9] used rheology to determine the morphology. Elasticity seems to be the key parameter. Indeed, the previous models show that the droplet relaxation caused by interfacial forces has strong impact on the elastic part of complex shear modulus. Interfaces bring additional elasticity. Then it is easily demonstrated that the interfacial area between the components is the cause of this additional elasticity. Figure 8.19 shows the storage modulus G of each component and of a PEO/PVdF-HFP blend (70/30), in the frequency range from 10−2 to 100 rad s. This figure depicts the influence of the morphology on G modulus, clearly visible at low frequencies. Comparison with the Kerner model demonstrates the excess of elasticity brought by the interface. At low compositions, the well-known behavior of droplets/matrix system is present as shown previously (the Palierne model). In the intermediate compositions, the shoulder of the G

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Storage modulus G' (Pa)

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15.8 rad.s−1

104

103

0.1 rad.s−1 Kerner

102 0.01 rad.s−1

101

1 102 10 ) 100 ad/s −1 r 10 ω( −2 y 10 c en 10−3 qu Fre 1

100

80 60 PEO 40 20 con tent (w% )

0

Figure 8.19 Storage modulus G versus frequency at 150◦ C for the complete composition range. (•) 0.01 rad.s−1 , (o) 0.1 rad.s−1 , (- - -) Kerner model prediction (Eq. (8.26)). Reprinted from [9]. Copyright (2004) with permission from Elsevier.

curve is less visible. Castro et al. [9] explained that the cause is the existence of complex morphologies, i.e., partially continuous and/or fully co-continuous morphologies. Few authors [3–5, 9] deal with the rheological behavior of the intermediate morphologies, in the composition range where the droplets/matrix is not available anymore. The superimposition of the continuity curve obtained by selective dissolution and the G evolution vs. composition shows that the two maxima observed on G agree well with the limits of a co-continuous zone. Authors [9] conclude that the aspect of the G curve is very similar to the evolution of interfacial area with composition measured by image analysis [4]. More general thermodynamical considerations show that the contribution to the module of the interface is directly linked to the interfacial area [86]. Considering a droplet of radius R, the variation of its surface after applying a strain ε is given by: A = A f − Ai = 4π π 2 (1 + 0.4ε2 ) − 4π π 2 ≈ π R 2 ε2

(8.31)

For N droplets in the system, the increase in free enthalpy G is: G = αN A ≈ αNπ R2 ε2 Knowing that for a volume fraction φ of the dispersed phase N = is given by: G ≈

α 2 φε R

(8.32) 3φ , 4π R 3

the free enthalpy per unit volume

(8.33)

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As G is close to the internal energy U for such system, it easily comes that:   Eε2 G = U = σ dε = Eεεd = 2

337

(8.34)

The extra modulus E brings by the interfacial increase is: E≈

α ϕ ≈ αQ R

(8.35)

where Q is the specific interfacial area. 8.4.3.2

Theoretical Description of the Low Frequency Viscoelastic Behavior of Blends

Although it was often observed experimentally on polymers blends, only one theoretical description of Figure 8.20 has been proposed [9], in order to link morphology and rheological dynamic data. In their approach, the morphology created during mixing is characterized by one parameter: the interfacial area. The rheological analysis is a tool to measure this area. The main assumption of their proposed model is that the elasticity in a blend of incompatible polymers is due to both components and also to the existence of the interfacial tension α which in connection with the interfacial area per unit volume Q defines the magnitude of the elastic forces. This is consistent with the work of Palierne [60], and Lee and Park [68] or Froehlich [69]. They consider as Elias et al. [66] that the storage modulus of such a blend can be written as the sum of a component involving the composition of the blend plus the interfacial area contribution. The composition effect is classically depicted by Kerner’s model. The effect of interfacial area is then simply written as G ∗α (ω, α) = kα Q

(8.36)

where k is a dimensionless constant. 104

100

Continuity (%)

80

G′(0.01 rad.s–1)

103 60

40 102 PEO

20

PVdF-HFP

Storage modulus G′ (Pa)

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20

40 60 80 PEO content (w%)

100

101

Figure 8.20 Overlay of co-continuity curves and storage modulus G vs PEO content ((•) co-continuity of PVdfHFP, (o) co-continuity of PEO, () G at 150◦ C and 0.01 rad.s−1 ). Reprinted from [9]. Copyright (2004) with permission from Elsevier.

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In their work they showed that the interfacial area is strongly dependent on the way the blend has been processed. They particularly focus on the competition between coalescence and dispersion during the blending stage. Thus, they proposed two regimes according to the composition of the blend. For low concentration of dispersed phase, rupture is the main mechanism of interface creation. Increase of the interfacial area by coalescence is considered as a second order mechanism. The interfacial area is then governed by the droplet size after break-up and consequently by the critical capillary number Ca and critical capillary number Cacrit . They found the following expression of the interfacial area: Q=3

φηm γ˙ αCacrit

(8.37)

where ηm is the viscosity of the matrix. Equation 8.37 shows that in the borders of the composition diagram, the area Q increases with the amount of dispersed component B. As soon as the volume fraction exceeds some percolation threshold, coalescence can no more be neglected in the establishment of the final morphology. Thus, in this intermediate region of concentration, Castro et al. [9] calculated the area Q by a balance between the dispersion and coalescence mechanisms as stated by Tokita [38]. They found that in this regime, the expression of Q is given by: Q=

ERφ π ηm γ˙ − 4Pα α

(8.38)

where P is probability of coalescence and ER is the macroscopic rupture energy. By contrast with the first regime, Q is clearly a decreasing function of φ when coalescence has occurred during the mixing stage. The maximum of the interfacial area is calculated at the intersection of the variations (37) and (38). Figure 8.21 shows that the model developed fits well with the evolution of the storage modulus of their PEO/PVdF-HFP binary system with only one adjustable parameter. However, the model predicts that the

Figure 8.21 Storage modulus G vs. frequency at 150◦ C and 0.01 rad.s−1 . (•) experimental data. (—) model prediction (Eqs. (37) and (38)). Reprinted from [9]. Copyright (2004) with permission from Elsevier.

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maximum of interfacial area which coincides with the limits of the co-continuous zone is independent of the interfacial tension. Anyway, this theoretical description allows the bringing together of mixing conditions, morphology and rheological characterization.

8.5 Microphase Separated Block Copolymers 8.5.1

Ordered State and Morphologies in Block Copolymers: The Case of SEBS Triblock

Microphase segregation in block copolymers is explained in terms of the volume fraction of the components and the segment–segment interaction parameter χ = (δ a − δ b )2 , where δ a and δ b are the solubility parameters of the respective chains: the inverse of χ reflects the degree of affinity between the segments [1]. The microphase-separated or ordered state is arranged typically in spherical, cylindrical (hexagonal), lamellar or bicontinuous double-diamond structures. Considering that the dimensions of the ordered structures or domains and the distances between them are of few tens of nanometers, these systems can be defined as self-assembled nanostructures. Since the interaction parameter χ varies linearly with the inverse temperature 1/T, increasing T leads to more affinity between the segments, resulting in a spatially homogeneous disordered system at sufficiently high temperatures. The commonly-named order–disorder transition (ODT) separates the two thermodynamic regimes. Among these complex polymer systems, styrenic block copolymers and in particular Poly (Styrene-bButadiene-b-Styrene) (SBS) triblock copolymers are industrially the most relevant and the best known scientifically. But in recent years Poly (Styrene-b-Ethylene-co–Butylene-b-Styrene) (SEBS) copolymers are progressively gaining ground. The technical applications of SBS and SEBS copolymers cover different areas, including the automotive, electronics and sports industries, as well as being used in civil engineering purposes such as road paving and waterproofing membranes. In many cases they are employed as plastics and bitumen modifiers to improve impact resistance at low temperatures. Gels for cables and other uses are also within the scope of styrenic TPEs, because they have a great oil absorption capacity, forming gels that can be applied to protect delicate pieces (optics or electronics). In particular, SEBS is intended for use where UV resistance, high service temperature, and processing stability are essential. As an example within the field of block copolymers, this chapter focuses on SEBS copolymers of different molecular weights which contain a polystyrene volume fraction of φ = 0.30. Transmission Electron Microscopy (TEM) results [70] have shown that these copolymers have a hexagonal arrangement of PS cylinder domains of in a Poly(Ethylene-co–Butylene) (PEB) matrix. In general, in microphase separated copolymers the ordered state is maintained in a wide range of temperatures, up to the order–disorder temperature (TODT ). For example, in the case of the SEBS copolymer with a PS volume fraction of φ = 0.30 an order–disorder temperature higher than 300◦ C has been detected [71]. In the disordered state the copolymer shows the typical viscoelastic response of a homogeneous polymer melt. Therefore, the respective ordered and disordered states are characterized by qualitative different lowfrequency dynamic viscoelastic features, as has been shown in the literature for different block copolymers [72–84]. Interesting efforts have been made to elucidate the relationship between morphology and rheology of ordered states, in particular to associate viscoelastic responses with respective cubic, hexagonal and lamellar phases [84]. Referring to the particular case of styrenic copolymers, a crucial difference between SBS and SEBS triblock copolymers is the ability of the latter to crystallize. In fact, either semi-crystalline or completely amorphous

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SEBS copolymers can be produced, depending on the ethylene/butylene proportion of the central PEB block which is included in the chemical structure as displayed below.

The effect of crystallization on the viscoleastic and mechanical properties of SEBS copolymers has been the subject of a number of papers [85–88].

8.6

Dynamic Viscoelastic Results of SEBS Copolymers

8.6.1 8.6.1.1

Low and Intermediate Frequency Viscoelastic Behavior Analysis of the Entanglement Plateau

The peculiar rheological behavior of block copolymers near the microphase separation temperature (that is to say, order–disorder temperature) was to a great extent clarified in 1984 by Bates [78]. It was observed that in the linear viscoelastic regime, at low frequencies, the storage (G ) and loss (G ) moduli of the microphase separated samples were strongly dependent on the phase state of these samples. This gives rise to rheologically complex materials, for which the time–temperature superposition (TTS) principle to obtain a reduced or master curve cannot be applied. Once microphase separation (self-assembled nanostructure) is eliminated, heating the block copolymer above its order–disorder temperature (TODT ), the sample satisfies TTS principles, behaving like a thermorheologically simple material. Certainly, complexity and simplicity concepts have provoked the interest of outstanding rheologists, like Markovitz [89] and Plazek [90], who conclude that thermorheological complexity can be proven, but simplicity cannot. This is relevant in the rheological analysis of block copolymers. Let’s consider, for instance, the results of a Poly (Styrene-bEthylene-co–Butylene-b-Styrene) (SEBS) for which the PS volume fraction is φ = 0.30 (poybutylene content 40%), the total molecular weight Mw = 75000 g/mol (distributed as 11250 g/mol-52500 g/mol-11250 g/mol) and the polydispersity index Mw/Mn = 1.2. The result of applying the time–temperature method to the dynamic viscoelastic data of this sample, obtained at different temperatures ranging from 90 to 270◦ C, is shown in Figure 8.22. In spite of the heterogeneous nature of the copolymer (separated microphases up to 300◦ C), reasonable master curves are obtained, in particular for the loss modulus G and the complex viscosity η* . This rather surprising result stands for the relativity of the concepts ‘thermorheologically simple’ and ‘thermorheologically complex’, which is reconsidered below. In any case, we remark that the rubbery or entanglement plateau characteristic of polymers is clearly envisaged. The comparison of our data with the entanglement moduli of PS (G 0N = 0.2 MPa), PE (G 0N = 2.6 MPa) and Polybutylene (PB) (G 0N = 0.2 MPa) reported in the literature [91], leads us to assume that the observed plateau should be associated with the entanglements of poly(Ethylene-co–Butylene) (PEB) polymer chains which constitute the matrix of the investigated SEBS copolymer. A necessary requirement for midblock PEB chains being entangled is that their Mw should be higher than the entanglement molecular weight, Me. This is fulfilled in the case of our sample, because PEB chains have a molecular weight (Mw = 52500 g/mol) much higher than the entanglement molecular weights of polyethylene and polybutylene, 1800 and 13000 g/mol [91], respectively. The entanglement modulus of the SEBS block copolymer here considered can be evaluated using the procedure

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7

10

6

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PE 6

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PS PB

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10

4

η' (Pa.s)

G' ; G" (Pa)

10

10

0

10

3

10

2

10 0,0001

-2

0,01

1

100

10

4

6

10

10

f (Hz)

Figure 8.22 Storage modulus (), loss modulus () and real part of the complex viscosity (x) as a function of frequency, for the SEBS sample described in the text. The shift factor aT is employed to superpose the data to a reference temperature T = 150◦ C. The lines indicate the respective values of the entanglement modulus of PS (G0N = 0.2 MPa), PE (G0N = 2.6 MPa) and Polybutylene (PB) (G0N = 0.2 MPa).

described by Ferry [92], consisting in integration over the appropriate maximum in G :

G 0N

2 = π

∞

G  (ω)dLnω

(8.39)

−∞

The fitting of the experimental data shown in Figure 8.23 to the loss modulus function G  = m 1 · em 2 (ln ω−m 3 )

2

(8.40)

and the corresponding integration in Eq. (8.39), lead to G 0N = 0.85 MPa, that can be taken as a valid reference for SEBS samples containing around 40% polybutylene, since the entanglement modulus is practically independent of the molecular weight. In view of the lower entanglement modulus value of PB (G 0N = 0.2 MPa), compared to that of PE (G 0N = 2.6 MPa), a G 0N decrease can be envisaged as polybutylene proportion in PEB is increased. 8.6.2

Thermorheological Complexity

The thermorheological complexity of the investigated SEBS sample is clearly noticed in tan δ vs G* plots, such as those shown Figure 8.24. This is owed to its self-assembled nanostructure, below the order-disorder transition. This rheological method, proposed by Mavridis and Shroff [93] to analyze the temperature dependence of dynamic viscoelastic data, has been revealed as the most efficient technique to investigate the thermorheological complexity of complex systems such as branched polymers and inmiscible polymer blends [94–99]. As an example of this assessment, we remark that whereas other rheological methods, like breaking the TSS principle in the building up of a master curve or log G vs log G plots [100], are not able to ascertain

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Figure 8.23

Loss modulus values of SEBS sample described in the text at T = 250◦ C fitted to Eq. (8.40).

40 270°C 250°C 230°C 210°C 190°C 170°C 150°C 130°C 110°C

35

30

δ

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20

15

10 4 10

5

6

10

10

7

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G*

Figure 8.24 Phase angle plot as a function of complex modulus for the SEBS sample described in the text. Failure of time–temperature superposition method is observed.

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thermorheological complexity of inmiscible PVC based blends, tan δ vs G* plots clearly account for a complex thermorheological response [94]. Besides the thermorheological complex behavior, in Figure 8.24 a phase angle maximum is observed at all temperatures. This result is interpreted as a consequence of the blocking effect of ordered PS cylinders on PEB chains motion, a subject clarified below.

8.6.3

Specific Mechanical Relaxation at Low Frequencies

As has been explained in a recent paper on the rheology of SEBS copolymers [70], the motion of each PEB chain as a whole is suppressed, because of the hindering effect of PS cylinders which constitute the hexagonal phase. This leads to a loss tangent decrease, as the frequency decreases (low values of G* in Figure 8.24). Actually, this hindering effect is better interpreted in terms of tan δ versus frequency results, shown in Figure 8.25. The existence of a low frequency tan δ maximum at a constant temperature has been reported in the literature for microphase-separated polystyrene-polyisoprene diblock copolymers (SI) [83,101] and Poly (Styrene-b-Isoprene-b-Styrene) triblock copolymers (SIS) [102, 103], but only the paper by Arevalillo et al. [70] offers an interpretation in terms of chain dynamics. Similarly to the case of polymer nanocomposites investigated by Fern´andez et al. [104], this mechanical relaxation represents the obstructing effect of a structure (ordered microdomains, in the case of block copolymers) to the mobility of the chain associated to flow. In the disordered state (T > TODT ) ordered microdomains disappear and chains are free to move: then tan δ → ∞ at low frequencies and the spectrum resembles that of homopolymers, with no secondary relaxation. Figure 8.26 shows the difference between the tan δ spectrum of a homogeneous copolymer melt (disordered state) and a block copolymer which contains ordered microdomains. The viscoelastic spectrum of an amorphous polymer is presented in Figure 8.26a (according to the description of Ferry), whereas Figure 8.26b is actually an adaptation of the viscoelastic pattern of ordered block copolymers, envisaged by Kossuth et al. [84]. 1 210°C 230°C 250°C 270°C

0,8

tan δ

0,6

0,4

0,2

0 0,01

0,1

1

10

f (Hz)

Figure 8.25 Mechanical relaxation (tan δ maximum) observed for the SEBS sample described in the text. Tan δ maximum shifts to higher frequency as temperature increases.

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tan δ

Flow

Flow

Glass Transition

Microdomain Relaxation

Glassy State

Entanglements

(a)

Glass Transition

Ordered Entanglements Microdomains

Glassy State ω

ω

(b)

Figure 8.26 The loss tangent as a function of frequency. (a) Homogeneous polymer melt; (b) microphase separated block copolymer. The data presented in Figure 8.25 correspond to the marked mechanical relaxation zone. Reprinted from [70]. Copyright (2008) with permission from Elsevier.

In homogeneous systems, as frequency decreases, the loss tangent passes through a minimum separating motions within entanglement strands and motions across entanglement loci. In the case of microphaseseparated copolymers (ordered microdomains at T < TODT) , the midblock polymer chains which form the matrix suffer a double constraint to motion due to: (a) Chain entanglements, provided that the molecular weight of the midblock polymer is higher than its entanglement molecular weight (Me); (b) A three-dimensional structure caused by ordered microdomains. According to this picture, the lowering tan δ values seen at high frequencies in Figure 8.25 should be associated with the entanglements of PEB chains (discussed above), whereas as frequency is decreased, the observed tan δ maximum marks a limit in the mobility of matrix chains in an ordered system. In view of the morphology of the SEBS copolymer here considered, what is detected through tan δ maximum is the blocking effect of PS cylinders on PEB chains motion. The characteristic time of the relaxation process is defined taking the inverse of the frequency at which (tan δ)max, occurs λ = 1/ωmax . At a given temperature, the chain mobility that requires times larger than λ (which corresponds to the low frequency zone with respect to the maximum in Figure 8.26b) is interfered. Assuming that the tan δ maximum observed in Figure 8.25 is associated with the mechanical relaxation depicted in Figure 8.26b, we can state that motions that imply distances larger than the strands between microdomains are blocked, since they require times t > λ. But entanglement slippage motions can be accomplished, because they are more local and they take place at shorter times, implying t < λ. As can be seen in Figure 8.25, tan δ maximum shifts to higher frequencies as temperature increases, which is equivalent to say that λ diminishes as temperature increases. This is not surprising, considering that the mobility of the polymer chains augments with temperature and, therefore, the blocking effect of ordered microdomains at T< TODT should be noticed at higher frequencies or shorter times. The dependence of the relaxation time with temperature follows an Arrhenius-like Eq. (8.41): Eλ

λ(T ) = A · e RT

(8.41)

where the activation energy is Eλ = 130 kJ/mole for the λ values obtained from Figure 8.25. Considering that the activation energies of flow of PE and PB are, respectively, Ea = 30 kJ/mole and Ea = 48 kJ/mole [105] and that Vinogradov and Malkin [106] give a value of Ea = 35 kJ/mole for ‘copolymers of ethylene with 1-butene’ (with no specification of ethylene/1-butene proportion), we deduce that the activation energy of the relaxation process in (tan δ)max , Eλ , is much higher than the activation energy of flow which accounts for the mobility of PEB chains. This indicates that λ reduction with temperature is not only due to

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an increase of mobility (viscosity reduction) of PEB chains. A progressive weakening of the anchoring effect of PS on PEB chains (at the interface PS/PEB) as temperature increases can be hypothetically considered to contribute also to λ decrease.

8.7 Flow-induced Morphological Changes 8.7.1

Order–Order Transition and Flow Alignment in Block Copolymers

The effect of flow fields on the morphology of block copolymers has attracted a handful of researchers over the last decades. The feature paper of Wiesner [101] on order and dynamics of block copolymers submitted to large amplitude oscillatory flow (LAOS) and the review paper of Hamley [107] on the use of strong shear fields, large extensional flows and large-amplitude oscillatory shear to align lamellar, hexagonal and cubic micellar morphologies of diblock and triblock copolymers, constitute relevant contributions to the topic. In the majority of the papers published so far alignment or mesophase order of morphologies is considered [108–112], disregarding the case of a flow field inducing a phase transition. The effect of extrusion flow to align the PS cylinders of a SBS copolymer was first studied by Keller et al. [113] in 1970. More recently, extrusion flows involving different shear rates have been used to orient the microphases of different multiblock polymers and, eventually, to destroy this orientation [111, 112]. Transmission Electron Microscopy (TEM) is often used to analyze orientation, but more conclusive results on oriented block copolymers are probably obtained with two dimensional (2D)-SAXS [114–116]. The order–order transition (OOT) (for instance from lamellar to hexagonal order or vice versa) induced by a flow field, by temperature or by interaction with another polymer has also been treated in the literature [117–123], although the results are not always conclusive. Kotaka et al. [124] report OOT in a SEBS copolymer induced by an elongational flow. Probably the most clear case of a shear-induced order–order transition (from spherical to hexagonal order), observed by in situ SANS, is shown by Koppi et al. [118] for a diblock copolymer. Indeed, detecting an order–order transition is a more difficult task than determining the order–disorder temperature, since the latter can be performed by differential scanning calorimetry (DSC) or more typically using dynamic viscoelastic results. To our knowledge, only one paper [120] reports rheological results of an OOT; in particular a hexagonal- to-cubic transition in a SIS triblock copolymer at 192◦ C. At this temperature an exothermic peak and a drop of two decades in the storage modulus G are observed. 8.7.2

Flow Alignment in a SEBS Copolymer

The effect of a continuous flow (constant shear rate γ˙ = 0.6 s−1 during 300 s) on the dynamic viscoelastic properties of the SEBS characterized above, is shown in Figure 8.27. Experimental details are given in reference [70]. Meaningful differences are found between the dynamic viscoelastic response of the nosheared sample and the sample which has been submitted to the continuous shear flow. Oscillatory (dynamic) flow measurements were performed immediately after cessation of flow, avoiding any recovery of the sheared sample. The behavior changes from an elastic dominant response (G > G ) (associated with the three-dimensional structure of ordered domains) observed for the no-sheared sample, to G > G noticed for the sheared sample. For the sheared sample a fall of more than one decade in G is observed at the lowest frequency. Significant results concern also the frequency dependence of the elastic moduli of the original and the sheared sample, in the lowest frequency region. The slopes, which correspond to the α values of the scaling law G ∝ ωα , are shown in Table 8.2 and compared with the α values given by Kossuth [84] for cubic, hexagonal and lamellar morphologies.

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G' Coulored symbols G" Empty symbols (Pa)

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6

10

5

10

4

10

3

No-sheared Sheared

10

2

0.01

0.1

1

10

100

ω (Hz)

Figure 8.27 Storage modulus G and loss modulus G of original (no sheared) SEBS sample and the same sample submitted to a shear flow. In the latter, oscillatory flow measurements were performed immediately after cessation of flow, to avoid any recovery (see Reference 70 for experimental details). The slopes of G at the lowest frequency are presented in Table 8.2. Reprinted from [70]. Copyright (2008) with permission from Elsevier.

An exclusively rheological understanding of the results can suggest an order–order transition (OOT), from a hexagonal to a lamellar phase, provoked by shear flow. This interpretation would be supported by the huge G decrease observed in Figure 8.27, which is concomitant with the aforementioned G results on OOT of SIS triblock copolymers [120], as well as by the coincidence of the α exponent (Table 8.2) of the sample submitted to flow with the exponent that would correspond to a lamellar order. However, TEM results presented in Figure 8.28 do not favor the hypothesis of an OOT being promoted during flow, but rather sustain that PS cylinders are aligned by flow, although an orientation rearrangement is noticed close to the edge. Therefore, flow alignment is at the origin of the differences, shown in Figure 8.27, between the dynamic viscoelastic response of the no-sheared sample and the sample which has been submitted to shear flow. As Table 8.2 α exponent of the scaling law G ∝ ωα according to Kossuth et al. [84] and fitting the SEBS results in the low-frequency region of Figure 8.27. Scaling laws G ∝ ω2 and G ∝ ω, with G > G , are observed for homogeneous melts, including block copolymers at T>TODT . Kossuth et al. [24]

SEBS sample

0 (cubic) 0.33 (hexagonal) 0.5 (lamelar)

0.25 (no sheared) 0.5 (sheared)

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Figure 8.28 TEM microphotographs of SEBS samples characterized in the text. (a) Sample quenched after compressing at 250◦ C. (b) Sample quenched after continuous shear flow (see [70] for experimental details). Flow direction is normal to the picture. Upper left part constitutes the edge of the sample, where a different morphology is observed. Reprinted from [70]. Copyright (2008) with permission from Elsevier.

PS cylinders are aligned in flow direction their blocking effect on PEB chains motion is reduced, which gives rise to a more frequency-dependent G (higher exponent in G ∝ ωα scaling law). This also explains the high dissipation level, associated with G > G observed at low frequencies, of the samples submitted to shear flow. The rheological consequences of flow alignment can also be interpreted in terms of the λ relaxation time defined taking the inverse of the frequency at which tan δ max is observed. As explained in the paper of Arevalillo et al. [70], an increase in relaxation time λ indicates that larger PEB chain strands are implicated in motion when PS cylinders are aligned by flow. The stability of the new morphological state, induced by flow, is investigated analyzing the evolution of the dynamic viscoelastic functions with resting time after flow is ceased. The effect of resting time on the elastic modulus G (ω) is shown in Figure 8.29: As resting time increases, the viscoelastic response tends to approach the behavior reported for no-sheared SEBS. The results of Figure 8.29 reveal that after flow is ceased, the cylinder’s orientation relaxes, although not completely. Therefore, a recovery from a flow-oriented cylindrical phase towards a non-oriented PS cylinders morphology can be deduced from frequency dependent rheological data, indicating that the sample tends to its equilibrium morphology in a no-sheared state. Our data show a saturation of the orientation recovery, after 60 min. This incomplete recovery indicates that at 250◦ C the initial equilibrium conditions are not reached after cessation of flow, which would suggest a certain recovery impediment caused at the interface with the rheometer plates.

8.8 Capillary Extrusion Rheometry Results of Block Copolymers 8.8.1

General Results of Styrenic Block Copolymers

Viscosity measurements carried out in capillary extrusion rheometers are very relevant in the field of applied rheology of polymers, because a wide range of shear rates (typically from γ˙ = 1−1 s−1 to more than 104 s−1 ),

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G' (Pa)

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10

No-sheared Sheared 5 min resting 10 min resting 30 min resting 45 min resting 60 min resting

3

10

2

10 0.01

0.1

1

10

100

ω (Hz)

Figure 8.29 Storage modulus vs. frequency for a SEBS sample (characterized in the text) at T = 250◦ C, measured at the indicated resting times after shear is ceased. Reprinted from [70]. Copyright (2008) with permission from Elsevier.

which covers the shear rates involved in compression molding, calendering, extrusion and injection molding, can be obtained. Moreover, measurements in a broad temperature interval can be performed. Therefore, the viscosity is measured in conditions close to industrial processing. Since 1965, the year styrene-based thermoplastic elastomers were presented by the Shell Chemical Company, the rheological studies of these microphase-separated block copolymers have been mainly devoted to dynamic viscoelastic measurements. Notwithstanding, a certain number of prominent works have dealt with capillary extrusion rheometry [111,125–132]. The influence of microphase separation on viscosity function η(γ˙ ) of SBS copolymers was reported by Holden et al. [125], remarking on the existence of two distinct viscosity–shear rate relationships. Depending on temperature, the transition from one state to the other takes place at different shear rates, although always at a shear stress of 105 Pa, as can be seen in Figure 8.30 for an SBS copolymer with a polystyrene matrix and spherical PB domains. This behavior is a consequence of the network structure which remains practically intact even at temperatures well above any of the Tg-s of the separated phases, preventing a complete dissipative flow and approaching the behavior of a solid. This aspect was pointed out by Vinogradov et al. [126], comparing the flow curves of a polybutadiene and an SBS (with polystyrene block inclusions in a PB matrix) copolymer of similar molecular weight. An adaptation of their results is shown in Figure 8.31 which resembles the situation observed upon addition of a more viscous polymer to a less viscous one, constituting an inmiscible blend [106]. Another interesting aspect of microphase-separated SBS copolymers, disclosed by Vinogradov et al. [126], is their viscoplastic and thixotropic behavior, that the authors identify with the response of conventional polymer melts with solid fillers. The results of the International Union of Pure and Applied Chemistry (IUPAC) investigation on SBS samples, developed by 14 research laboratories and reported in 1980 [127], confirm the thixotropy detected at low shear stresses.

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106 150°C

VISCOSITY, polse

150°C 175°C 105 200°C

SH

EA

R

ST

RE

SS

104

=1

06

dy ne

s/c

m2

103

1

102

10

SHEAR RATE,

103

104

sec–1

Figure 8.30 Viscosities of SBS copolymer with 39% styrene content. Reprinted from [125]. Copyright (1969) with permission from John Wiley & Sons.

More recently, the capillary rheometer results of a linear and a star-shaped SBS copolymer, both of similar molecular weight, were compared [112], observing that the star-shaped sample was less shear thinning, as accounted by a higher value of the ‘α’ index of the Cross model [134]: η = η0 (1 + (γ˙ τ0 )a )

(8.42)

Where η0 is the Newtonian or linear (shear rate independent) viscosity and τ 0 a relaxation time associated with disentanglement time of entangled chains. Lingae-Jorgensen [128] summarized the basic features which determine the correlation microstructurerheology-processing for block copoymers, remarking the effect of the difference δ a − δ b between the solubility

log σ (Pa)

6

4

2

0 –6

–4

–2

0 . Log γ (s–1)

2

4

Figure 8.31 Flow curves of polybutadiene (black circles) and SBS (white circles) of similar molecular weight. Reprinted from [126].

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105

VISCOSITY (Pa Sec)

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104

103

102

101 4 10

S-EB-S I

S-EB-S II

200 °C 220 °C 240 °C 260 °C 280 °C 300 °C

200 °C 220 °C 240 °C 260 °C 280 °C 300 °C

105

106

105 SHEAR STRESS (Pa)

106

105

106

Figure 8.32 Viscosity as a function of shear stress at the indicated temperatures for an SBS and two SEBS samples of total molecular weight Mw = 190,000 g/mol (SEBS I) and Mw = 70,000 g/mol (SEBS II) and PS volume fraction of φ = 0.30. Reprinted from [129]. Copyright (1987) Hanser Publishers.

parameters of the respective phases. It is observed that processability improves as this difference is reduced. This is on the basis of the viscosity differences between SBS and SEBS copolymers, remarked by Gergen et al. [129] and shown in Figure 8.32. The viscous behavior of SBS is typical of a well-developed flow in a thermoplastic polymer, whereas SEBS, which possesses a bigger interaction parameter, χ = (δa − δb )2 , shows an absence of dissipative (viscous) process, similar to crosslinked rubbers.

8.8.2

Viscosity and Flow Instabilities in SEBS Copolymers

The capillary rheometer results of the viscosity function η(γ˙ ) of four SEBS samples of PS volume fraction φ = 0.30 and molecular weights Mw ranging from 50,000 g/mol to 200,000 g/mol, are shown in Figure 8.33. The viscosity curves of the lowest molecular weight sample (Mw = 55000 g/mol) respond to the classical shear thinning behavior of thermoplastic melts (which fit to the Cross model [134], Eq. (8.42), marking an incipient linear (shear rate independent) viscosity at low shear rates. The sample which possesses a total molecular weight Mw = 75000 g/mol also shows a Cross-like behavior, but only at the highest temperature T = 250◦ C. Therefore, we can assert that, depending on molecular weight and temperature, SEBS copolymers can flow as typically conventional polymer do, notwithstanding that in many cases the linear viscosity region is absent and shear thinning extends to the whole shear rate range investigated. The latter behavior is noticed in Figure 8.33a for the samples of molecular weights Mw = 135000 g/mol and 185000 g/mol. In particular, the highest molecular sample is characterized by a power law dependency η ∝ γ˙ −1 (slope −1 in log η –log γ˙ plots) which denotes a pseudoplasticity index n = 0 according to the power law for the shear stress

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η

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100 10

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. -1 10 4 γ (s ) (d)

Figure 8.33 Viscosity as a function of shear rate at 190◦ C, 210◦ C, 230◦ C and 250◦ C for SEBS samples of PS volume fraction φ = 0.30, polybutylene content 40% and molecular weights: (a) Mw = 55000 g/mol (825038500-8250). Data are fitted to the Cross model (Eq. (8.42)). (b) Mw = 75000 g/mol (11250-52500-11250). (c) Mw = 130000 g/mol (19500-91000-19500). (d) Mw = 185000 g/mol (27750-129500-2775).

σ ∝ γ˙ n . A zero value for the pseudoplasticity index signifies that the velocity profile in the capillary is flat, giving rise to the so called ‘plug flow’, which has been remarked in some SEBS samples by Ghosh et al. [133]. Moreover, unlike for the lowest molecular weight sample (Figure 8.33a), no effect of temperature on viscosity is observed in the case of SEBS-s of Mw molecular weights Mw = 135000 g/mol and 185000 g/mol. Therefore, an activation energy Ea = 0 in Eq. (8.43): Ea

η(T ) = B · e RT

(8.43)

is deduced for high molecular weight SEBS samples, leading us to consider that these block copolymers slide in the capillary wall, as a crosslinked rubber or a solid of a relatively low friction coefficient Actually, as stated by Vinogradov and Malkin [106], when the melt slips in the capillary wall the concept of the activation energy flow becomes inapplicable, because viscous flow itself disappears and a plastic deformation occurs by a deactivation mechanism. When clearly an actual capillary flow occurs, as is the case of the SEBS sample having Mw = 55000 g/mol (Figure 8.33a), the activation energy decreases with shear rate, varying from Ea = 58 kJ/mol at 15 s−1 to Ea = 18 kJ/mole at 1500 s−1, which denotes that different flow units are involved depending on shear rate (thermorheological complexity). As stated above,

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the activation energy of flow of midblock PEB chains of SEBS samples should be close to 35 kJ/mol. For its part, Ghosh et al. [133] investigated capillary flow of SEBS samples of polystyrene volume fraction φ = 0.30 in a Monsanto Processability Tester, concluding that the activation energy of the highest molecular weight sample (Mw = 163333 g/mole) was zero, because plug flow and wall slippage occurred. For a sample of molecular weight Mw = 102222 g/mole these authors report an activation energy which decreases from 16 kJ/mole to 2.4 kJ/mole, when the shear rate is changed from 125 s−1 to 630 s−1 , and in the case of a sample of Mw = 58889 a value Ea = 7.5 kJ/mole is given, practically independent of shear rate. Considering that the lowest reported Ea value for polymers is 20 kJ/mole, for polysiloxane [106], we have to be cautious with the values given by Ghosh et al., which may have been obtained under no actual shear flow conditions. One of the difficulties encountered in the processing of block copolymers in general and SEBS triblocks in particular is the occurrence of flow instabilities leading to extrudate distortions at shear rates typically involved in polymer processing. Flow instabilities constitute a topic recurrently studied by polymer scientist over the past 50 years. Different types of extrudate distortions (matt or absence of specular gloss, sharkskin, slip-stick, waviness, ripple, gross melt fracture, spurt) have been reported in many papers and outstanding reviews have been written about the matter [135–139], notwithstanding the fact that few papers refer to instabilities of block copolymers. So far, the role microdomains may have on the assumed flow instabilities of flexible midblock chains, like polybutadiene and PEB, has not been disclosed. Phatak et al. [110] analyze die extrusion instabilities of lamellae forming triblock and pentablock copolymers composed of poly(cyclohexylethylene) and polyethylene, observing that sharkskin is much more severe in the pentablock copolymer, which is attributed to a shear induced disordering near the die wall in contrast with lamellae parallel alignment in triblock sample. In the case of styrene-isoprene-styrene (SIS) triblock copolymers, Ferri and Canetti [131] observe that the critical shear stress for the inception of slip-stick and spurt phenomena increases with increasing polystyrene content. The striking phenomenon of flow split [140], which is a very severe instability leading to the scission of a polymer melt into two branches at the exit of a capillary die, reported for the first time in ethane/propylene copolymers, has been observed for SEBS copolymers [80, 130, 132]. As reported by Santanach et al. [130], SEBS copolymers forming hexagonal-packed cylinders of PS give flow split, but SEBS samples forming ordered spherical microphases do not develop this severe instability. The authors verify that only for the latter SEBS the loss modulus G overtakes the storage modulus G , which is a symptom of less melt elasticity. This supports the hypothesis, advanced by Fernandez et al. [140], of flow split being due to a high elasticity of the molten polymer. In Figure 8.34 some representative microphotographs of extrudates of the SEBS samples described in Figure 8.33 are shown. Actually, smooth and practically transparent extrudates are only obtained with the lowest molecular weight sample, Mw = 55000 g/mol, at shear rates up to 72 s−1 at 250◦ C and below this shear rate at lower temperatures. The latter sample at higher shear rates, as well as the rest of SEBS samples contemplated in Figure 8.33, show either matt (no gloss) or severe sharkskin or flow split at any shear rate and temperature. It is worth pointing out that the less severe kind of distortion (smooth although without specular gloss, matt, extrudates) is systematically observed for the highest molecular weight sample (Mw = 186000 g/mol). This is certainly due to the aforementioned plug flow which, in turn, arises as a consequence of the very big melt elasticity of this highly entangled sample. Sharkskin and its ultimate consequence, flow split, are related to an entanglement–disentanglement process, as that proposed by Wang [137], but highly entangled polymers show slip at the wall, since no disentanglement takes place during flow. The effect of both restricting phenomena, anchoring of PEB chains into PS domains and entanglements, on the viscoelastic terminal or flow zone of SEBS samples of different molecular weights is shown in Figure 8.35. The influence of the molecular weight on tan δ versus frequency plots is interpreted considering the general picture depicted in Figure 8.26. Due to the flattening and spreading out of the entanglement minimum as

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(b)

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(d)

Figure 8.34 Microphotographs of SEBS extrudates obtained at 250◦ C. (a) Smooth; Mw = 55000 g/mol at γ˙ = 3 s−1 . (b) Matt; Mw = 186000 g/mol at γ˙ = 80 s−1 . (c) Severe sharkskin or melt fracture; Mw = 75000 g/mol at γ˙ = 3 s−1 . (d) Flow split; Mw = 75000 g/mol at γ˙ = 80 s−1 .

the molecular weight increases, the relaxation associated with the hindering effect of PS cylinders on PEB chains mobility is shifted to lower frequencies (out of the experimental range). A consequence of this feature is the very high melt elasticity (very low tan δ value in the whole frequency range) of the highest molecular weight SEBS (Mw = 185000 g/mol), associated with the inability of PEB to disentangle, which provokes the observed slip at the wall and plug-flow in extrusion measurements.

1,2

1

55000 75000 130000 185000

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0 0,01

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Figure 8.35 Loss tangent as a function of frequency at T = 250◦ C for SEBS samples. The molecular weights of the samples (in g/mol units) are given in the inlet.

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8.9

Summary

Multiphase polymeric systems, such as immiscible blends and microphase-separated block copolymers, exhibit specific rheological behavior. In the case of multiphase blends, this peculiar viscoelastic response originates from the microscopic heterogeneity described by the presence of local discontinuities along with interfacial forces. Consequently, the shape relaxation of the dispersed phase generates extra stress only encountered for such systems. The particular viscoelastic behavior of block copolymers arises as a consequence of a physical network where elastic chains are held together by glassy or crystalline microdomains. This deep link between rheology and morphology becomes a powerful tool to characterize polymer blends as well as block copolymers. On the other hand, rheological conditions during mixing originate from the final morphology of blends and, for its part, thermo-mechanical conditions greatly affect the domain organization of block copolymers. Microrheology governs the basic mechanism of multiphase blends such as break-up and coalescence. The fine understanding of the impact of rheology on morphology and vice versa is then the key to control, analyze and to tailor the blends to the final application. Many theoretical or phenomenological models have attempted to capture the macroscopic behavior of polymers blends considering the whole system as a homogeneous one. Starting for microscopic relaxation mechanisms, these models have turned out to be in agreement with the experimental data in linear viscoelastic conditions. However, even if the interface elasticity is well captured by the theories, all of them fail when the concentration of the dispersed phase increases or when the morphology differs from classical droplet/matrix as for the co-continuous systems. The analysis of the restrictions imposed by styrenic microdomains to the motion of elastic PEB chains in SEBS triblock copolymers, is taken as a reference for a general vision on the rheology of block copolymers. As dynamic viscoelastic experiments reveal, local motions like entanglement slippage are permitted for midblock chains, but motions that imply distances larger than strands between microdomains are blocked. The application of a continuous flow leads to an alignments of the microdomains causing a significant alteration of the viscoelastic response, since the blocking effect is reduced. Nonlinear rheological experiments, such as those carried out in a capillary rheometer in a wide range of shear rates, disclose a liaison between flow instabilities and viscoelastic response at low frequencies.

References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13.

Bates, S.F., Science, 898–904, 251 (1991). Galloway, J.A., Koester, K.J., Paasch, B.J., Macosko C.W. Polymer 423–8, 45 (2004). Huitric, J., Mederic, P., Moan, M., Jarrin, J., Polymer 4849–56, 39 (1998). Galloway, J.A., Macosko, C.W., Polym. Eng. Sci., 714–727, 44 (2004). Steinmann, S., Gronski, W., Friedrich, C., Polymer 6619–6629, 42 (2001). Potschke, P., Paul, D.R., J Macromol Sci, Polym Rev., 87–141, C43(1) (2003). Bourry, D., Favis, B.D., J Polym Sci., 1889–1899, 36 (1998). Steinmann, S., Gronski, W., Friedrich, C., Rheol Acta 77–86, 41 (2002). Castro, M., Carrot, C., Prochazka, F., Polymer, 4095–4104, 45 (2004). Avgeropoulos, G.N., Weissert, F.C., Bidisson, P.H., B¨ohm, G.C.A., Rubber Chem. Technol., 93, 49 (1976). Karger Kocsis, J., Kallo, A., Kuleznev, V.N., Polymer, 279, 25 (1984). Rumscheidt, F.D., Mason, S.G., J. Colloid. Sci., 238, 16 (1961). Grace, H.P., Eng. Found. Res. Conference Mixing 3rd N. H. Andover (1971) republished as Dispersion Phenomena in High Viscosity Immiscible Fluids Systems And Application of Statics Mixers as dispersion Mixers in Such Systems. Chem. Eng. Commun., 225, 14 (1982). 14. Wu, S., Polym. Eng. Sci., 335, 27 (1987).

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Rheology and Viscoelasticity of Multiphase Polymer Systems: Blends and Block Copolymers 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57. 58. 59. 60. 61. 62.

355

Favis, B.D., Chalifoux, J.P., Polym. Eng. Sci., 1591, 27 (1987). Taylor, G.I., Proc R Soc., 138, 41 (1932). Taylor, G.I., Proc R Soc., 146, 501 (1934). Van Oene, H.J., Colloid. Inter. Sci., 448, 40 (1972). Elmendorp, J.J., Maalcke, R.J., Polym. Eng. Sci., 1041, 25 (1985). Levitt, L., Macosko, C.W., Pearson, L.C. D., Polym. Eng. Sci., 1647, 36 (1996). Min, K., White, J.L., Fellers, J.F., J. Appl. Polym. Sci., 2117, 29 (1984). Vinogradov, G.V., Krasnikova, N.P., Dreval, V.E., Kotova, E.V., Plotnikova, E.P., Pelzbauer, Z., Inter. J. Polym. Mater., 187, 9 (1982). Jansen, J.M.H., Meijer, H.E.H., J Rheol., 597–608, 37 (1993). de Bruijn, R.A., PhD Thesis, Eindhoven University (1989). Karam, H.J., Bellinger, J.C., Ind. Eng. Chem. Fundam., 576, 7 (1968). Tavgac, T., Ph.D Thesis, University of Houston, Texas (1972). Rayleigh, J.W.S., Proc. London. Math Soc., 10, 1 (1879). Tomotika, S., Proc. Royal Soc., London, 322, A150 (1950). Stone, H.A., Bentley, B.J., Leal, L.G, J. Fluid Mech., 131, 173 (1986). Janssen, J.M.H., PhD Thesis, Eindhoven Univ. Technol., The Netherlands (1993). Minale, M., Mewis, J., Moldenaers, P., AICHE J., 943–50, 44 (1998). Minale, M., Moldenaers, P., Mewis, J., Macromolecules, 5470–75, 30 (1997). Chesters, A.K., Trans. IchemE Sci., 259–70, A 69 (1991). Ottino, J.M., DeRoussel, P., Hansen, S., Khakhar, D.V., Adv. Chem. Eng., 105–204, 25 (1999). Tucker III, C.L., Moldenaers, P., Annu. Rev. Fluid Mech., 177–210, 34 (2002). Ablazova, T.I., Tsebrenko, M.B., Yudin, A.B.V., Vinogradov, G.V., Yarlykov, B.V., J. Appl. Polym. Sci., 1781, 18 (1975). Krause, S., Macromolecules, 1602, 13 (1980). Glasstone, S., Laidler, K.L., Eyring, H., Theory of Rate Processes, McGraw-Hill, New York (1941). Bingham, E.C., Fluidity and Plasticity, page 86, McGraw Hill, New York (1922). Friedmann, E.M., Proter, R.S., Tranc. Soc. Rheol., 493, 19 (1975). Doi, M., Edwards, S.F., J. Polym. Sci., Polym. Lett. Ed., 265, 19 (1981). Cassagnau, Ph., Montfort J.P., Marin, G., Monge, Ph., Rheol. Acta., 156–167, 32 (1993). Tsenoglou, C., Macromolecules, 1761–1767, 24 (1991). Majest´e, J.C., Allal, A., Montfort, J.P., Rheol. Acta, 477–482, 42 (2003). Pathak, J.A., Van-Gurp, M., J Appl. Polym. Sci. 1245–1249, 78(6) (2000). Zhiyi,Y., Han Chang, D., Macromolecules, 2104–2118, 41(6) (2008). Einstein, A., Ann Phys, 289–306, 19 (1906). Simha, R., J. Appl. Phys., 1020, 23 (1952). Utracki, L.A., Polymers Blends and Alloys, Hanser Publishers, Munich (1989). Mooney, M.J., Colloid Sci., 162, 6 (1951). Krieger, I.M., Dougherty, T. J., Trans. Soc. Rheol., 137–152, 3 (1959). Kitano, T., Kataoka, T., Rheol Acta, 390, 20 (1981). Kitano, T., Kataoka, T., Nagatsuka, Y., Rheol Acta, 20, 23 (1984). Agoudjil, B., Ibos, L., Majest´e, J.C., Candau, Y., Mamunya, Ye.P., Composites Part A: Applied Science and Manufacturing, 342–351, 39 (2008). See, H., Jiang, P., Phan Tien N., Rheol. Acta, 131–7, 39 (2000). McEwan, M., Green, D., Soft Matter, 1705–1716, 5 (2009). Oldroyd, J.G., Proc R Soc., 122–32, A218 (1953). Kerner, E.H., Proc Phys Soc., 808–13, A69 (1956). Lewis, T., Nielsen, L., J Appl Poly Sci., 1449–71, 14 (1970). Palierne, J.F., Rheol. Acta, 204–214, 29 (1991). Bousmina, M., Rheol Acta, 73–83, 38 (1999). Graebling, D., Muller, R., Palierne, J.F., Macromolecules, 320, 26 (1993).

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356 63. 64. 65. 66. 67. 68. 69. 70. 71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81. 82. 83. 84. 85. 86. 87. 88. 89. 90. 91. 92. 93. 94. 95. 96. 97. 98. 99. 100. 101. 102. 103. 104. 105. 106. 107.

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Handbook of Multiphase Polymer Systems Friedrich, C., Gleinser, W., Korat, E., Maier, D., Weese, J Rheol., 1411–1425, 11, 39 (6) (1995). Elias, L., Fenouillot, F., Majeste, J.C., Cassagnau, P., Polymer, 6029–6040, 48 (2007). Vermant, J., Cioccolo, G., Golapan Nair, K., Moldenaers, P., Rheol Acta, 529–538, 43 (2004). Elias, L., Fenouillot, F., Majeste, J.C., Martin, G., Cassagnau, J. Polym. Sci.: Part B: Polymer Physics, 1976–1983, 46 (2008). Bousmina, M., Rheol Acta, 251–254, 38 (1999). Lee, H.M., Park, O.O., J Rheol., 1405–25, 38 (1994). Froehlich, H., Sack, R., Proc R Soc., 415–30, A185 (1946). Arevalillo, A., Mu˜noz, M.E., Santamaria, A., Fraga, L. and Barrio, J.A., European Polymer J., 3213–3221, 44 (2008). Kim, J.K., Jung, D.S. and Kim, J., Polymer, 4613–4624, 34 (1993). Chung, C.I. and Gale, JC., J Polym Sci., 1149–1156, 14 (1976). Gouinlock, E.V. and Porter, R.S., Polym Eng Sci., 534–542, 17 (1977). Chung C.I. and Lin M.I., J Polym Sci., Polym Phys., 545–553, 16 (1978). Cohen, R.E. and Ramos, A.R., Macromolecules, 131–134, 12 (1979). Widmaier, J.M. and Meyer, G.C., J Polym Sci., Polym Phys., 2217–2225, 18 (1980). Cohen, R.E. and Wilfong, D.E., Macromolecules, 370–375, 15 (1982). Bates, F.S. Macromolecules, 2607–2613, 17 (1984). Han, C.D. and Kim, J., J Polym Sci, Polym Phys., 1741–1764, 25 (1987). Han, C.D., Kim, J. and Kim, J.K., Macromolecules, 383–394, 22 (1989). Rosedale, J.H., Bates, F.S., Macromolecules, 2329–2338, 23 (1990). Bates, F.S., Rosedale, J.H., Fredrickson, G.H., J Chem Phys., 6255–6270, 92 (1990). Zhang, Y., Wiesner, U., Macromol. Chem. Phys., 1771–1784, 199 (1998). Kossuth, M.B., Morse, D.C., Bates, F.S., Journal of Rheology, 167–196, 43 (1999). Morton, M., in Thermoplastic Elastomers, (Chapter 4), N.R. Legge, G. Holden, H.E. Schroeder (Eds), Hanser Publishers, Munich, Vienna, New York (1987). Stadler, R., Gronski W., Colloid Polym. Sci., 323–331, 264 (1986). Sierra, C.A., Galan, C., Fatou, J.G., Parellada, M.D., Barrio J.A., Polymer, 4325–4335, 38 (1997). Indukuri, K.K., Lesser, A.J., Polymer, 7218–7229, 46 (2005). Markovitz, H.J., J Polym. Sci. Symp., 50, 431–456 (1975). Plazek, D., Journal of Rheology, 40, 987–1014 (1996). Graessley, W.W., Polymeric Liquid Networks: Dynamics and Rheology, Taylor and Francis Group, London and New York (2008). Ferry, J.D., Viscoelastic Properties of Polymers, John Wiley & Sons Inc., New York (1980). Mavridis, H., Shroff, R.N., Polym. Eng. Sci., 1778–1793, 32 (1992). Z´arraga, A., Pe˜na, J.J., Mu˜noz, M.E., Santamaria, A., J Polym. Sci.: Part B: Polym. Phys., 469–477, 38 (2000). Trinkle, S., Walter, P., Friedrich, C., Rheologica Acta, 103–110, 41 (2002). Vega, J.F., Fernandez, M., Santamar´ıa, A., Mu˜noz-Escalona, A., Lafuente, P., Macromol. Chem. and Physics, 2257–2268, 200 (1999). Gurp, M.V., Palmen, J., Rheo. Bulletin, 67, 5–8 (1998). Gonz´alez, O. Pe˜na, J.J., Mu˜noz, M.E., Santamaria, A., P´erez-Lepe, A., Mart´ınez-Boza, F., Gallegos, C., Energy and Fuels, 1256–1263, 16 (2002). Rojo, E., Pe˜na, B. Mu˜noz, M.E., Santamaria, A., Macromol. Chem. Phys., 1781–1788, 207 (2006). Han, C.D., Lem, K., Polym. Eng. Rev., 135–142, 2 (1982). Wiesner, U., Macromol. Chem. Phys., 3319–3352, 198 (1997). Ryu, C.Y., Lee, M.S., Hajduk, D.A. and Lodge, T.P., J. Polym Sci.: Polym Phys., 2811–2823, 35 (1997). Brown, K., Hooker, J.C. and Creton, C., Macromol. Mater. Eng., 163–179, 287 (2002). Fern´andez, I., Santamaria, A. Mu˜noz, M.E., Castell, P., Euro. Polym. J., 3171–3176, 43 (2007). Porter, R.S., Johnson, J.F., J Polym. Sci.: Part C, 373–380, 15 (1966). Vinogradov, G.V., Malkin, A. Ya., Rheology of Polymers, Mir Publishers Moscow-Springer-Verlag Berlin (1980). Hamley, I.W., J Phys.: Condensed Matter, 643–671, 13 (2001).

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108. Mortensen, K., Theunissen, E., Kleppinger, R., Almdal, K., Reynaers, H., Macromolecules, 7773–7781, 35 (2002). 109. Sebastian, J.M., Lai, C., Graessley, W.W., Register, R.A., Marchand, G.R., Macromolecules, 2700–2706, 35 (2002) 110. Phatak, A., Macosko, C.W., Bates, F.S., Hahn, S.F., Journal of Rheology, 197–214, 49 (2005). 111. Leist, H., Geiger, K., Wiesner, U., Macromolecules, 1315–1317, 32 (1999). 112. Geiger, K., Knoll, K., Langela, M. Rheol. Acta, 345–355, 41 (2002). 113. Keller, A., Pedemonte, E., Wilmouth, F.M., Nature, 538–539, 225 (1970). 114. Maring, D., Wiessner, U., Macromolecules, 660–662, 30 (1997). 115. Langela, M., Wiesner, U., Spiess, W. and Wilhelm, M., Macromolecules, 3198–3204, 35 (2002). 116. Oelschlaeger, C., Gutman, J.S., Wolkenhauer, M. Spiess, H.W., Knoll, K., Wilhelm, M., Macromol. Chem. Phys., 1719–1729, 208 (2007). 117. Park, C., Simmons, S., Fetters, L.J., Hsiao, B. Yeh, F., Thomas, E.L., Polymer, 2971–2977, 41 (2000). 118. Koppi, K.A., Tirrell, M., Bates, F.S., Amdal, K., Mortensen, K., Journal of Rheology, 999–1027, 38 (1994). 119. Heck, B., Arends, P., Ganter, M., Kressler, J., St¨uhn B., Macromolecules, 4559–4566, 30 (1997). 120. Kim, J.K., Lee, H.H., Gu, Q.J., Chang, T., Jeong, Y.H., Macromolecules, 4045–4048, 31 (1998). 121. Krishnamoorti, R., Modi, M.A., Tse, M.F., Wang, H.C., Macromolecules, 3810–3817, 33 (2000). 122. Jeong, U., Lee, H.H., Yang, H., Kim, J.K., Okamoto, S., Aida, S., Sakurai, S., Macromolecules, 1685–1693, 36 (2003). 123. Sakurai, S., Okamoto, S. and Sakurai, K., in Developments in Block Copolymer Science and Technology, I.W. Hamley (Ed.), John Wiley and Sons, Chichester, UK (2004). 124. Kotaka, T., Okamoto, M., Kojima, A., Kwon, Y.K., Nojima, S., Polymer, 3223–3231, 42 (2001). 125. Holden, G., Bishop, E.T., Legge, N.R., J Polym. Sci.: Part C, 37–57, 26 (1969). 126. Vinogradov, G.V., Dreval, V.E., Malkin, A.Ya., Yanovsky, Yu.G., Barancheeva, V.V., Borisenkova, E.K., Zabugina, M.P., Plotnikova, E.P. and Sabsai, O.Yu., Rheol. Acta, 258–263, 17 (1978). 127. Ghijsels, A., Raadsen J., Pure and Applied Chem. Report 3159, 52 (1980) IUPAC Report. 128. Lingaae-Jorgensen, J., in Processing, Structure and Properties of Block Copolymers (Chapter 3), M.J. Folkes (Ed.), Elsevier Applied Science Publishers, London and New York (1985). 129. Gergen, W.P., Lutz, R.G. and Davison, S., in Thermoplastic Elastomers (Chapter 14), N.R. Legge, G. Holden, H.E. Schroeder (Eds), Hanser Publishers, Munich, Vienna, New York (1987). 130. Santanach Carreras, E., El Kissi, N., Piau, J.M., J. Non-Newtonian Fluid Mechanics, 1–21, 131 (2005). 131. Ferri, D., Canetti, M., Journal of Rheology, 611–624, 50 (2006). 132. Piau, J.M., Santanach-Carreras, E., Suarez Dabo, E., J. Non-Newtonian Fluid Mechanics, 139–148, 292 (2007). 133. Ghosh, S., Bhowmick, A.K., Roychowdhury, N., Holden, G., J Appl. Polym. Sci., 1621–1628, 77 (2000). 134. Cross, M.M., J Appl. Polym. Sci., 765–769, 13 (1969). 135. Tordella, J.P., in Rheology: Theory and Applications (Chapter 2), F.R. Eirich (Ed.), Academic Press, New York and London (1969). 136. Vinogradov, G.V., Protasov, V.P., Dreval, V.E., Rheol. Acta, 46–61, 23 (1984). 137. Wang, S.-Q., Advances in Polymer Science, 227–275, 138 (1999). 138. Denn, M.M., Annu. Rev. Fluid Mechanics, 265–287, 33 (2001). 139. Hatzikiriakos, S.G., Migler, K.B. (Eds), Polymer Processing Instabilities, Marcel Dekker, New York (2005). 140. Fern´andez, M., Santamaria, A., Mu˜noz-Escalona, A., Mendez, L., Journal of Rheology, 595–602, 45 (2001).

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9 Thermal Analysis of Multiphase Polymer Systems Gy¨orgy J. Marosi Budapest University of Technology and Economics, Budapest, Hungary

Alfr´ed Menyh´ard Budapest University of Technology and Economics Budapest Hungary Hungarian Academy of Sciences, Chemical Research Centre, Budapest, Hungary

G´eza Regdon Jr. University of Szeged, Szeged, Hungary

J´ozsef Varga Budapest University of Technology and Economics Budapest Hungary Hungarian Academy of Sciences, Chemical Research Centre, Budapest, Hungary

9.1 Introduction Thermal analysis, starting from early works on DTA, has a long history, which was always in close relationship with material science and technology [1]. All the aspects of thermal analysis have been thoroughly reviewed in several hundreds of pages of the relevant papers and books [2, 3]. The aim of this chapter has to be limited to discussing only those selected techniques that are considered particularly important for multiphase systems. This limitation, however, gave a freedom for discussing the methods in an order that differs from the conventional sequence of subjects. The structure of this chapter is organized in a way that leads the reader from the direct visual evaluation to the more complex methods. Therefore the survey starts with the most spectacular thermooptical microscopy (TOM) coupled with the polarized light microscopy (PLM), which is rarely discussed among the other thermal analytical methods in spite of the fact that it provides direct and sensitive information about the structure formation and phase transitions under different thermal Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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conditions. The most widely-used conventional differential scanning calorimetry (DSC) is then discussed in detail, followed by an introduction to the more subtle modulated DSC methods. Localized and microchip techniques are introduced briefly just in order to indicate the direction of the development, and finally a selection from the wide variety of coupled thermogravimetric methods is given.

9.2

Thermooptical Microscopy

The visible route of thermal analysis is thermooptical microscopy (TOM). Optical micrographs, recorded with digital camera and suitable software during the run of a temperature program, promote the understanding of the development of the structure under different thermal conditions including thermal transitions. TOM using polarized light microscope is especially useful for studying polymer systems of different polymorphic forms. The optical character (the sign of the birefringence) of morphological entities can be determined by using lambda-plate located diagonally between polarizers [4]. Since the refraction index is higher in the chain direction than perpendicular to the chains, TOM-PLM studies performed using lambda plate give information about the orientation of the polymer chains. Furthermore, size distribution, occasional agglomeration of the additives and phase diagrams of the two-component polymer systems can be determined by TOM. The use of the twin slices technique (neat polymer and its modified version attached to each other) provides a simultaneous comparison under completely identical thermal conditions. Isotactic polypropylene (iPP) and its multicomponent systems are the objects investigated most widely by TOM [5–8 and references therein] because its alpha and beta polymorphic forms can be distinguished clearly [9, 10]. It was observed, for instance, that tensile stress arising on the surface of vacuum bubbles formed at the late stage of the crystallization of iPP induces beta-phase [11]. One of the most spectacular phenomenon taking place during the crystallization of the polymers is the formation of transcrystalline structure (TCS) on the surface of substances having nucleating ability [12–20]. A scheme of the transcrystallization is depicted in Figure 9.1. The modification of the crystalline structure is a possibility to influence the properties of a crystalline polymer relatively wide range according to the requirements of the application field. Nucleating agents (NAs) are used in industrial practice in order to modify the crystalline structure of the polymers. NAs are introduced into the polymers usually in 100–10000 ppm and they are present in the polymer matrix as finely-dispersed heterogeneous particles [17]. Some of the recently-used NAs can dissolve in the polymer melt resulting in a more homogeneous distribution of the additive [21–24]. The influence of a highly efficient alpha-nucleating

A

B

C

D

E

F

G

Figure 9.1 Scheme of the transcrystalline growth induced by heterogeneous nuclei formed at different times on the surface of an extraneous substance. Notes: Nuclei A, B, C, E and G are present at t = 0, D is formed at t = 0.5 and F at t = 1 time unit. Spherulit segments formed with time lag (D and F) are included by those started earlier. Contour lines indicate the growth front at different time [13].

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a tc = 3 min

b tc = 15 min

c tc = 40 min

d tc = 60 min

361

Figure 9.2 Isothermal crystallization of a twin iPP slice consisting of a non-nucleated (lower) and nucleated (upper) part in the presence of a highly efficient alpha-nucleating agent (NA21) at Tc = 140 ◦ C, (tc = 3, 15, 40 and 60 min). Nucleator concentration: 2000 ppm.

agent on the crystallization and on the structure of iPP is shown in Figure 9.2 using a twin slice. The nucleated (lower) part of the slice crystallizes at a high rate yielding microcrystalline structure, while the non-nucleated (upper) part begins with a long time lag and big spherulites are formed. It can be also seen that a transcrystalline-like crystal front is growing from the surface of the crystallized part of the slice. N,N -dicyclohexyl-2,6-naphthalene-dicarboxamide (NJS), an efficient nucleating agent, is soluble in iPP melt [21]. Its influence on the formation of the supermolecular structure of iPP is shown in Figure 9.3. This series of micrographs demonstrate the recrystallization of the soluble nucleating agent in the form of needle crystals (Figure 9.3(a)) and the growth of transcrystalline layers on their lateral surface at Tc = 135 ◦ C (Figure 9.3(b) and (c)). The transcrystalline layer consists of the alpha-phase with lower and the beta-phase with higher birefringence. The structure remaining after separate melting of the beta-form with lower melting point is seen in Figure 9.3(d). Sorbitol derivatives introduced in iPP in order to improve its optical properties are usually soluble in iPP melt [23, 24]. They exert an alpha-nucleating effect if their concentration exceeds a threshold value. Some pigments act as alpha-nucleating agents as well. The influence of a red pigment of strong alpha-nucleating effect, is demonstrated in Figure 9.4. Based on these micrographs, it can be stated that some pigments exert nucleating effects commensurate of that of highly efficient nucleating agents. It is well known that some fillers, such as talc, possess an alpha-nucleating ability for iPP [25–27]. Formation of a transcrystalline layer on the surface of talc is reported [27]. Surface treatment of the fillers can strongly modify their nucleating ability. Calcium carbonate, wollastonite and mica, for instance, treated with

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a Tc=135, tc = 0 min

b tc = 5min

c tc=30 min

d T=156

Figure 9.3 Effect of a soluble beta-nucleating agent (N,N’-dicyclohexyl-2,6-naphthalene-dicarboxamide) on the crystallization of iPP at Tc = 135 ◦ C (tc = 0, 5 and 30 min) and melting of partially-crystallized slice at T = 156 ◦ C. [21]. NA concentration: 500 ppm.

pimelic acid (a component of the calcium stearate and pimelic acid two-component beta-nucleating agent), possess moderate beta-nucleating ability [28, 29]. On the surface of fibers with strong nucleating ability, a well-developed transcrystalline structure can be formed as depicted in Figure 9.5 by a carbon fiber-reinforced iPP [14]. Some other organic fibers with moderate or weak nucleating ability (PET, PAN, polyamides, aramides [13] cellulose and other natural fibers [15]) induce the formation of a row of individual spherulites on the surface of the fibers as shown in the case of PET fiber in melt of iPP (Figure 9.6) [16]. Lowering Tc proliferate of the number of spherulites formed on the surface of PET fibre during the stepwise crystallization of iPP matrix (Figure 9.6(b) and (c)). Several fibers possess negligible or no influence on the crystallization of the iPP matrix. For example, glass fiber, being amorphous, is an inert additive in iPP [13]. In polymer blends, most additive polymers are incompatible and form heterogeneous structures. Several additive polymers (polyethylene, PVDF, etc.), including liquid crystalline ones [19], possess alphanucleating ability. The growth of a transcrystalline layer on their surface has been also observed by TOM (Figure 9.7) [20]. The effect of compatibilizer in polymer blends is well observable by TOM. Beyond the reduction of the size of dispersed particles the compatibilizer weakens or suppresses the alpha-nucleating effect of PA-6 in beta-nucleated blend of iPP and PA-6 (30).

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b tc = 12 min

c tc = 3 min

d tc = 180 min

363

Figure 9.4 Isothermal crystallization of an iPP twin slice consisting of a non-pigmented (left side) and a pigmented part (right side) and containing a red pigment (Cromophthal Red 2028) at Tc = 135 ◦ C (tc = 2, 12, 30 and 180 min).

Figure 9.5 Formation of a transcrystalline layer of iPP on the surface of the carbon fiber at Tc = 135 ◦ C.

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a Tc=135 °C, tc=30 min

b Tc=125 °C, tc=38 min

c Tc=125 °C, tc=40 min

Figure 9.6 Stepwise crystallization of iPP on PET fiber. (a) – Tc1 = 134 ◦ C, tc = 30 min; (b) – Tc2 = 124, tc = 38 min; (c) – Tc2 = 124, tc = 40 min [16].

iPP

iPP

T=150 °C

PVDF

T=130 °C

PVDF

Figure 9.7 Transcrystallization of iPP on the surface of crystalline PVDF [20].

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9.3 Differential Scanning Calorimetry Differential scanning calorimetry (DSC) is a tool for studying the thermal transitions quantitatively. Melting and crystallization characteristics, glass transition, chemical reaction, degradation behavior and even the specific heat capacity can be monitored by a single DSC measurement. All experimental techniques are based on the comparative measurement of thermal events in a specimen and in a reference material. The detailed introduction of the principles of DSC techniques is beyond the scope of this book. Several books are dedicated to collecting and introducing useful knowledge relating to the principles of thermal analysis of polymeric materials [31–35]. The two basic methods are the heat flux DSC and power compensated DSC. The former measures the temperature difference (T) between the sample and the reference material (the heat flux is proportional with the absorbed or released heat of the sample but it has to be corrected by the heating rate imbalance, the asymmetry of thermal resistance and capacitance of the sample versus reference platforms), The power compensated DSC method consists of two independently controlled furnaces, in which the absorbed heat (endothermic processes) and released heat (exothermic processes) are compensated by heating the sample or the reference cell respectively. The heat flux DSC is able to study the thermal transition in relatively large samples; the sample mass could even be 100 mg. Moreover, the application of this technique is advantageous in the study of gassing specimens. Power compensation DSC is flexible, because the measuring cell is small – sample mass is usually around 5–10 mg. This technique is available to study rapid thermal transitions, because the small measuring cells have low time constants [31, 34, 35]. Although, the basic principles of calorimetry have not changed during the last five decades, the development of experimental techniques is attractive as a consequence of continuously expanding computational and microelectronic possibilities. Versatile calorimetric techniques focus on the monitoring of thermal transitions at fast scanning rates [36–39]. Such approaches include the HyperDSC [37, 38] and Rapid-Heating-Cooling DSC (RHC-DSC) [39] techniques, which allow quantitative analysis up to heating rates of 750 and 2000 ◦ C/min respectively. The rapid heating and cooling make possible to investigate the thermal transitions under conditions representing certain industrial practices or accidents. Moreover, the fast scanning rate calorimetry (FSC) creates extended possibilities for studying time-dependent transitions. Recently-developing techniques are the stepwise isothermal segregation technique (SIST) and successive self-nucleation and annealing (SSA), which provide information about the fine molecular structure of crystalline polymers and copolymers [40, 41]. These calorimetric methods are based on the subsequent melting of the samples after controlled crystallization. Usually the sample is cooled above their crystallization temperature and held there for long periods (20 to 50 min). Subsequently the sample is cooled by 5–10 ◦ C with a high cooling rate and held there again isothermally for long time. These steps are repeated several times till the whole sample crystallizes. The characteristic shape of the subsequent melting peak will provide information about the molecular structure. The distribution of co-monomer in a copolymer and even the tacticity distribution in a stereoregular semicrystalline polymer can be estimated based on these experiments. The term ‘multiphase polymer system’ covers a wide variety of materials. The simplest multiphase polymer is the crystalline polymer itself, because it consists of both a crystalline and an amorphous phase. The thermal transitions of both phases can appear as signals on the calorimetric trace. The DSC trace recorded during heating of poly(ethylene-terephthalate) (PET) is a good example for introducing the thermal transitions in the amorphous and crystalline phases in one polymer (Figure 9.8). The low flexibility of PET chains results in a hindered tendency for crystallization and therefore PET can be cooled below its glass transition temperature (Tg ) without crystallization [42]. The calorimetric curve of an amorphous PET sample fabricated from the neck of a PET bottle is presented in Figure 9.8. Change of segmental mobility of amorphous PET is indicated by an endothermic step of the specific heat capacity (cp ) curve, the inflection point of which is attributed conventionally to Tg . The glass transition of highly crystalline

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Heat flow (mW) Endo

Exo

cold crystallization

Tmf

Tg

Tmp

0

50

100

150

200

250

300

Temperature (°C) Figure 9.8 The DSC trace of PET registered with heating rate of 10 ◦ C/min.

polymers is usually not observable in conventional DSC signals. As can be seen in Figure 9.8, an endothermic peak is superimposed on the stepwise change of cp , because the heating rate is faster than the relaxation rate of the segments in the vicinity of Tg . The exothermic peak around 150 ◦ C relates to the crystallization of amorphous PET below its melting point (Tm ). The crystallization process occurring during heating between Tg and Tm is called cold crystallization. The endothermic enthalpy change (Hm ) with peak temperature of 250 ◦ C (Tmp ) is related to melting of the crystalline phase of PET. The heating rate (vh ) influences the DSC signals as shown in Figure 9.9. At higher heating rates the step representing glass transition increases (facilitating its detection) and shifts toward higher temperatures. The cold crystallization, being dependent on segmental mobility, shifts to higher temperatures as well.

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0

50

100

150

200

250

300

Temperature (°C) Figure 9.9 The melting curves of PET recorded at heating rates of 10, 20 and 40 ◦ C/min respectively.

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The melting peak covers a broad temperature range, which represents crystallites with different lamellar thicknesses and degree of perfection [43]. The degree of crystallization (X) can be estimated according to Eq. (9.1). X=

Hm Hm0

(9.1)

Hm0 is the equilibrium enthalpy of fusion, which is available in the literature or can be determined by different extrapolation methods, for example by extrapolation of melting enthalpies of isothermally crystallized samples to Tm0 [44–47]. However, we have to point out here that Hm depends on the experimental condition and literature values of Hm0 usually have large scattering due to the way of its determination [8]. Therefore, the crystallinity values calculated according to Eq. (9.1) should be treated carefully. The shape of the melting curve (melting profile) reflects the distribution of lamella thickness, which can be determined this way [43]. The lamella thickness determines the melting temperature according to the Gibbs-Thomson equation:   2σe (9.2) Tm = Tm0 1 − Hv0  where Tm is the actual melting temperature of the polymer, Tm0 is the equilibrium temperature of melting, σ e is the free energy of the folded surface of a lamella, Hv0 is the equilibrium enthalpy fusion and  is the thickness of the lamella. A more concrete description of multiphase polymer systems is multicomponent polymer systems such as polymer blends and composites. The thermal transitions of all phases appear on the calorimetric trace individually. The heat flow changes of each process are superimposed onto each other if more processes are taking place simultaneously. The thermal events registered on the DSC traces relate only those phases where they are taking place respectively. Therefore, the composition of the specimen should be known in order to achieve reliable quantitative results. If the composition is unknown it should be determined by other techniques (for example TGA). Figure 9.10 demonstrates the enthalpy of fusion of PP-based composites containing Na-montmorillonite (NaMMT) as filler.

95 90

Enthalpy of fusion (J/g)

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85 80 75 70 Experimentally measured ΔHm Corrected ΔHmc by mass ratio

65 0

3

6

9

12

15

Filler content (mass%) Figure 9.10

Enthalpy of fusion values in PP/NaMMT composite as a function of composition.

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The small increase in Hm at low filler content is the consequence of the moderate nucleating effect of NaMMT. The Hm then decreases rapidly with increasing filler content, because the mass ratio of the polymer phase is decreasing. A considerably different tendency of Hm can be observed if the experimentally registered Hm is corrected by the mass ratio of the composite (Hmc ). Hmc is constant in a relatively broad composition range (1–10 mass%) and starts to decrease above 10 mass% of filler due to the hindered mobility of polymer fraction in the presence of a large amount of NaMMT. The crystallization of the polymer shifts towards the higher temperatures in the presence of nucleating agents (NA). Fairgrieve gives a summary about the recently-used nucleating agents in industrial practice [48]. The efficiency of an NA can be characterized by the shift of peak temperature of crystallization (Tcp ). Thierry et al. [49] proposed an efficiency coefficient in order to quantify the nucleating efficiency (Eq. (9.3)): E=

100(Tcp − Tcp0 ) (Tcp,max − Tcp0 )

(9.3)

where T cp0 and T cp,max are the crystallization peak temperatures of the non-nucleated polymer and the self seeded polymer respectively. This scale presents a reliable and quantitative value for efficiency. Fillon et al. [40] provide efficiency factors for several soluble organic NAs in iPP. The main drawback of this method is the determination of T cp,max , which depends on the final temperature of heating above Tm . The efficiency of an NA can be characterized by the shift of Tcp in its presence (Tcp ). This shift is presented for several NAs in Figure 9.11. NAs 1–3 are heterogeneous nucleating agents. In their presence Tcp increases monotonously. However, it reaches saturation above a certain concentration of NA. Slightly above the saturation point is the optimal NA content from the point of view of efficiency. In the presence of soluble clarifiers (NA 4 and 5) efficiency has a threshold concentration. Larger amounts of these NAs should be introduced than their solubility limits, because in other cases they can not act as heterogeneous nucleator. (The conditions of cooling determines the recrystallization of a soluble NA from the iPP melt [49, 21]. The nucleating effect can be explained by the ‘matching lattice size’ theory developed by Alcazar et al. [50]. In multicomponent polymer blends those phases could act as nucleator, which crystallizes at higher temperatures than the other components. The iPP/poly(vinylidene-fluoride) (PVDF) blend is a good example

Peak temperature of crystallization, Tcp (°C)

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118 116 114 112 110 108 106

NA 1 NA 2 NA 3 NA 4 NA 5

104 102 100

0

500 1000 1500 2000 2500 Nucleating agent content (ppm)

Figure 9.11 Peak temperature of crystallization as a funtion of nucleating agent content in the presence of different nucleators in iPP.

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Half time of crystallization, t1/2 (min)

Thermal Analysis of Multiphase Polymer Systems 120

120

100

118 116

80

114

60

112

40

110

20 0

108 0

5 10 15 20 PVDF content (wt%)

106 25

369

Peak temperature of crystallization, Tcp (°C)

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Figure 9.12 The half time and the peak temperature of crystallization in an iPP/PVDF blend as a function of PVDF content. The t1/2 values are obtained during isothermal experiments at Tc = 135 ◦ C [51].

for blends, where the additive polymer (PVDF) has strong nucleating effect on iPP [20, 51] as shown in Figure 9.12. The Tmp of the two components does not change considerably, which proves that there is no significant interaction between the completely immiscible iPP and PVDF and both components crystallize individually. The kinetics of nucleation and crystallization can be evaluated based on isothermal and non-isothermal experiments [32–35]. The half time of crystallization (t1/2 ) decreases dramatically in the presence of PVDF (see Figure 9.12). Recent studies discussed that the distribution of a nucleating agent depends on the polarity of the phases. iPP/Polyamide 6 (PA6) blend are good examples, because the polar nucleating agent is encapsulated in the polar PA6 phase [20, 30]. The distribution of the nucleating agent can be influenced by the modification of the interphase between the components by adding a compatibilizing agent to the blend [30]. The presence of a compatibilizer influences the nucleating effect of the polymer phases as well [52, 53]. The strength of interaction between two crystalline polymers can be characterized on the basis of the depression of melting temperature by the Flory-Huggins interaction parameter. Nishi and Wang [54] adapted the Flory-Huggins theory, to determine the interaction in crystalline polymer blends. The interaction parameter can be obtained by Equation (9.4): 1 RV2 1 − 0 =− χ12 (1 − ϕ2 )2 Tm Tm H2 V1

(9.4)

where Tm and Tm0 are the experimental and the equilibrium melting temperatures, R is the universal gas constant, V is the molar volume, H is the enthalpy of fusion and ϕ is the volume fraction of the components. The indexes of 1 and 2 represent the two components respectively. χ 12 is the Flory-Huggins interaction parameter. Equation (9.4) can be rearranged into the following form (Eq. (9.5)) by the introduction of the density of interaction energy (B): 1 ϕ1



1 1 − 0 Tm Tm

 =−

χ12 =

BV2 ϕ1 H2 Tm

BV1 RT

(9.5)

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Slope: 0.1125 2 R : 0.987

3

0.30 0.28

0

(1/Tm-1/T m)*10

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0.2

0.4

0.6

0.8

1.0

2

φ1

Figure 9.13

The Nishi-Wang plots of iPP/PP-g-MAH blend.

Both Eqs (9.4) and (9.5) are available to determine χ 12 . One technique is to plot the experimental results according Eq. (9.4), which should result in a straight line. The interaction parameter can be deduced from the slope [55–59]. The Nishi–Wang plots of iPP/iPP grafted by maleic acid (PP-g-MAH) is demonstrated in Figure 9.13, where PP-g-MAH is expected to interact with iPP [60]. (PP-g-MAH is mainly used as a compatibilizing agent in heterogeneous polymer blends based on iPP.) Three different types of PP-g-MAH polymers have been investigated in a recent work [60]. It was pointed out that the compatibilizing efficiency of PP-g-MAH depends predominantly on the strength of interaction between iPP and PP-g-MAH. In strict thermodynamical terms, polymer blends are miscible if they form one single phase system at molecular level. The miscibility of polymer blends can be characterized by the glass transition temperature of the components. Figure 9.14 shows schematically the Tg as a function of composition in a polymer blend.

Glass transition temperature, Tg (°C)

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a Tg (2)

b

c

b

a Tg (1) 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0

Composition ratio Figure 9.14

The schematic diagram of Tg as a function of the composition of polymer blends.

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Tg (1) and Tg (2) are the glass transition temperatures of the components respectively. Curve c, changing linearly with the composition ratio of the components, symbolizes the case of complete miscibility. Cases a and b represent complete immiscibility and partial miscibility respectively. A necessary, but not sufficient, condition of miscibility is that the free energy of mixing should be negative (Gmix < 0); however, we have to notice that the mixing entropy of polymers is usually too small to result in negative Gmix . Therefore, most of the polymer blends are only partially miscible or completely immiscible. The Tg of a miscible polymer blend can be predicted by the Fox equation (9.6) [60, 61]. 1 W1 W2 = + Tg Tg1 Tg2

(9.6)

where W 1 and W 2 are the mass fractions of the components and Tg , T g1 and T g2 are the glass transition temperatures of the blend and the two components respectively. In the case of partially miscible polymer blends the Flory-Huggins interaction parameter (χ 12 ) between the phases can be calculated on the basis of the shift of Tg of the components. The interaction parameter can be easily calculated by the technique suggested by Kim and Burns [62]. The free energy of mixing (Gm ) of a partially miscible polymer blend can be expressed as follows on the basis of the Flory-Huggins theory: G m = n 1 ln φ1 + n 2 ln φ2 + χ12 φ1 φ2 (n 1 m 1 + n 2 m 2 ) RT

(9.7)

where n1 and n2 are the molar amounts and φ 1 and φ 2 are the volume fraction of component 1 and 2 respectively. The m1 and m2 are essentially the degree of polymerization relating to the molar volumes of the components. M n,1 m1 ρ1 = m2 M n,2 ρ2

(9.8)

M n,i is the number average molar mass of the ith component and ρ i is the density of components 1 and 2 respectively (i = 1 or 2). The choice of the lattice site volume here can be arbitrary (usually the larger one is chosen), but once a site volume has been chosen for one of the components, it must be the same for the other components as well. The chemical potential of mixing of the components can be obtained as the partial derivative of Equation (9.9) with respect to ni .  μ1 = ln φ1 + 1 − RT  μ2 = ln φ2 + 1 − RT

m1 m2 m2 m1

 φ2 + m 1 χ12 φ22 

(9.9) φ1 +

m 2 χ12 φ12

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Glass transition temperature, Tg

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PVC/PMMA PVC/PS PMMA/PS

PVC PVC PMMA

PMMA PS PS

120 110 100 90 80 70 0.0

0.2

0.4

0.6

0.8

1.0

Volume fraction Figure 9.15

The Tg transition of components in PMMA/PVC blends determined by DSC at 20 ◦ C/min.

Denoting the two conjugal phases by single and double primes at equilibrium state, we have μ1 = μ1 and μ2 = μ2 . Thus Eq. (9.10) results in the interaction parameter:

χ12 =

χ12 =

  φ1 m1 (φ2 − φ2 ) ln  + 1 − φ1 m2 m 1 (φ2 2 − φ2 2 )   φ  m1 (φ1 − φ1 ) ln 2 + 1 − φ2 m2

(9.10)

m 1 (φ1 2 − φ1 2 )

With this simple method the interaction parameter can be deduced based on the Tg data. However, it should be pointed out that Tg determined by thermoanalytical techniques depends strongly on the experimental conditions and consequently the obtained interaction parameter depends on the reliability of the experiments. The Kim and Burns model [62] usually provides small positive numbers. The smaller this number the stronger the interaction between the phases. Figure 9.15 represents the change of Tg in several amorphous polymer blends [63]. A steep increase of Tg can be seen in the case of poly(vinyl-chloride (PVC)/poly(methylmethacrylate) (PMMA) blends. The average χ 12 parameter is ∼1.8* 10−3 , which indicates relatively strong interaction between PVC and PMMA. It has to be pointed out that only one Tg is present if the volume fraction of PVC is smaller than 0.7, which indicates that PVC and PMMA are miscible polymer pairs. The shift of Tg in PMMA/polystyrene (PS) and PVC/PS blends is much smaller than that in the previously discussed PVC/PMMA blends and two separate Tg can be observed in the two latter blends in almost the whole concentration range. The presence of two Tg in the whole concentration range hints that PVC and PMMA are only partially miscible with PS. Despite the partial miscibility the shift of Tg indicates interaction between the polymers. The average χ 12 in PVC/PS blends is in the range of 2.5–3.0* 10−3 , which also indicates quite strong interaction, but weaker than that of the miscible PVC/PMMA blend. It should be noted that the main drawback of the method of Kim and Burns is that the shift of Tg is small in most cases. Therefore, the scattering of the determined χ 12 values is relatively large. The other questionable point is

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that the choice of the lattice size volume is arbitrary, but it has strong influence on the estimated interaction parameter. Thus the interaction parameter calculated according to the Nishi–Wang or Kim–Burns models based on melting temperature or glass transition temperature respectively depends always on the method of calculation. Therefore, only those parameters are comparable which were obtained according to the same method. Nevertheless, the presented methods are effective and useful tools to compare the interaction between the phases in different multicomponent systems.

9.4 Temperature Modulated Differential Scanning Calorimetry In the case of conventional differential scanning calorimetry (DSC), a linear heating signal is applied to a sample, while the temperature and energy associated with the response are measured. The thermal characterization of polymeric materials is relatively complicated, because all properties have pronounced time dependency as well. Therefore, the characteristic values determined by the DSC technique are representative not only to the physical constants but also to the experimental conditions and techniques. In order to separate complex transitions into more easily interpretable curves, temperature-modulated DSC was introducted in the early 1990s. This type of dynamic measuring model involved the innovation of software modification techniques to allow the superimposition of a modulation on the underlying signal, as described in details in relevant reviews [64–67]. Also different manufacturers use different terminology for a variety of types of modulation and output, all such techniques are here called temperature modulated DSC (TMDSC), or modulated-temperature DSC (MTDSC); (alternating differential scanning calorimetry, or ADSC, is also widely used). In most cases the temperature is varied in a sinusoidal manner as a function of time and is superimposed on the average heating rate (Figure 9.15). The phase shift (delay) that occurs between the perturbation and the response can be deconvoluted mathematically (by means of Fourier analysis). The deconvolution yields the ‘total’ signal, which is close to the standard DSC, the ‘reversing’ quantity, which is the effect of the modulation alone and the non-reversing one, which represents the kinetic component of the heat flow. Specific heat changes are always visible in the reversing DSC curve, while the time-dependent processes, such as relaxation; re-crystallization; curing; decomposition; or evaporation, are always apparent in the non-reversing DSC curve. The method determines the Cp in the quasi-isothermal mode. Accuracy of Cp determination is reduced at higher modulation frequencies (due to thermal diffusivity within the sample). Blank and calibration measurements (using a heat capacity calibration standard) will eliminate cell asymmetry and will enhance the accuracy of Cp measurements even at higher modulation frequencies. In order to keep the signal damping low the mass of the sample should be as small as possible. For the same reasons high frequencies are not used. (The signal damping is compensated mathematically [68].) Normally the cycle time is not less than 1 minute. The average heating rate lies between –5 K/min and +5 K/min (the maximum heating and cooling rates of the method must be lower than the maximum heating and cooling rates of the modules). Amplitude of 1 K is normally used. TMDSC is advantageously used in the case of overlapping effects. Glass transitions can be separated well from decomposition, relaxation, loss of adsorbed water, or cold-crystallization processes, as it allows both temperature-dependent and time-dependent processes to be separated. For facilitating this capability, a new advanced multi-frequency technique has been recently developed [69]. In contrast to the earlier single frequency methods, which overlaid the isothermal temperature or heating ramp with a (usually) sinusoidal temperature modulation of just one frequency, the multi-frequency temperature modulation technique uses a large number of frequencies [70]. The basic idea of this concept is to overlay the isothermal or ramped temperature with a time series of stochastic (random) temperature pulses of different durations. As a result, the underlying heating rate is modulated with a broad band of frequencies. Sample characteristics can be determined this way as a function of time and temperature over a wide frequency range.

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Figure 9.16

Sinusoidal (a) and random (b) pulse temperature profiles.

The measurement, using temperature pulses, and the subsequent evaluation process can be seen schematically in Figure 9.16. The primary data of multi-frequency measurements are evaluated according to the steps described in Figure 9.17. The advantages of multi-frequency TMDSC methods are as follows:

r r r r r r

frequency-dependent effects (e.g., glass transitions) can be very easily distinguished from frequency independent effects; simultaneous measurement of sample properties over a large frequency range can be performed (instrumental influences on the time and temperature function can be avoided); low energy transitions at close-lying temperature can be identified (high sensitivity and high resolution); the pulse response provides a very accurate way for determination of the quasi-static (frequencyindependent) heat capacity (cp ); accurate frequency-dependent heat capacity values can be determined based on one single measurement; heat capacities can be determined even if the effects overlap.

If the temperature modulation is sufficiently small, one can assume that the current state of the sample is almost unaffected and that it is in equilibrium. One can therefore describe the sample in a limited temperature range as a linear system.

Figure 9.17

The evaluation process of multi-frequency TMDSC.

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From the measured heat flow Ømeasured (t,T) both the reversing heat flow Øreversing (t,T) and the non-reversing heat flow Ønon-reversing (t,T) can be calculated. The total heat flow is the sum of Øreversing (t,T) and Ønon-reversing (t,T) according to the equation / total (t, T ) = O /reversing (t, T ) + O /non-reversing (t, T ) O where: heat flow: Ø (t,T), reversing component: Øreversing (t,T) = mcp0 (t,T) ß, non-reversing component: Ønon-reversing (t,T), mass: m, specific heat capacity: cp (t,T), heating rate: ß. Multi-frequency TMDSC provides information about the frequency dependence of glass transition relaxation process of polymers being determined by the molecular dynamics and predictions can be made regarding the stability of the material’s structure. Biocomposites are examples of the advantageous use of TMDSC. It is widely used to examine not only basic and active ingredients (monocomponent systems) of pharmaceuticals but also mixtures, multiparticle systems and even products of food industry or biological tissues [71–75]. The DSC curves of pharmaceutical formulations often exhibit several overlapping thermal effects. The total heat flow curve of a hydrophilic polymer matrix mostly shows a broad vaporization peak that overlaps the peaks of glass transition and frequency independent phase transition. The quasi-static heat capacity curve helps to identify these peaks, while the evaluation of the non-reversing heat flow curve yields the amount of moisture loss [76]. In the case of thermoset polymers vitrification occurs during curing when the glass transition temperature (Tg ) of the reaction mixture exceeds the curing temperature. A nonlinear relationship was found between the apparent vitrification time and log(frequency) [69]. In thermoset composite materials of high fiber content, the glass transition is usually broad and is accompanied by only a small change in Cp. This change can be masked by overlapping post-curing reaction and onset of decomposition. Conventional DSC therefore shows only an exothermic post-curing reaction followed by decomposition. Using the TMDSC method, the glass transition of the material can be clearly seen in the reversing heat flow curve, while the non-reversing heat flow shows the post-curing reaction and the onset of decomposition (Figure 9.18). The future for modulated temperature DSC experiments lies with decreased sample volume because only measurements made on small samples can follow the temperature modulation perfectly. This is one of the reasons why the micro-thermal measurements became important.

Figure 9.18

TMDSC monitoring of the curing of epoxy resin.

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Figure 9.19

9.5

July 6, 2011

Schematic drawing of (a) Wollaston probe and (b) ‘bow-tie’ probe of SThM.

Micro- and Nanothermal Analysis

TMDSC gives quantitative information on the formation of interfaces in multiphase polymer systems, and allows the corresponding shift in glass transition temperature to be detected. It is adequate when spatial discrimination is not required. However, the detection of the local changes of the thermomechanical characteristics at the phase borders is crucial to the understanding of the macroscopic properties (e.g., debonding, delamination processes). Such analysis is possible using micro-thermal analysis. Thermal analysis on a micro scale can be performed using an atomic force microscope (AFM), in which the original tip is replaced by a small resistive heater, thus forming a so-called ‘scanning thermal microscop’ (SThM) [77–79]. One type of this resistive heater is a thin, V-shaped noble metal wire, called a ‘Wollaston probe’ (Figure 9.19(a)), while another ‘bow-tie’ probe is formed by patterning an abrupt taper in a palladium wire that runs across the flattened apex of a 2.5 × 2.5 μm square (Figure 9.19b). These probes can serve for heating the tip and for determining the local temperature characteristics with about 1–0.1 μm spatial resolution. The various types of local thermal analyses are performed by ramping the temperature of the tip at a selected areas using direct (dc thermal), or modulated temperature (ac thermal amplitude/phase) programs. By using a reference probe and modulated programs the micro equivalent of a TMDSC experiment can be performed on a small volume (a few cubic micrometres) of material. However, due to unknown mass, the method is quantitative only in terms of identifying transition temperatures, but not quantitative in measuring transition enthalpies (L-TMDTA). Temperature oscillation is about one degree amplitude and the frequency is usually set at 1 kHz (higher than the bandwidth of the feedback). The chief motive for using temperature modulation is to look into the sample to different depths (which can be controlled by the frequency), providing a tomographic three-dimensional image. Heat transfer from the tip to the specimen results in local thermal expansion of different values, which can be detected using the AFM z-axis feedback system. For performing localized versions of thermomechanical analysis (L-TMA), the probe is first pressed against a chosen region of the sample, leading to a deflection of the cantilever, with a given initial force. The temperature-calibrated probe is then heated so as to apply an upward and downward temperature ramp, and its deflection along the z-axis (perpendicular to the sample surface) is recorded. If the material undergoes a phase transition, its mechanical properties vary (softening), so that the tip indents the sample and the cantilever deflection changes. In multiphase polymer systems the gradients in heat transfer properties across an interphase (IP) and spatial variations in local softening or glass transition temperature (Tg ) can be determined by thermal imaging and LTMA respectively. These analyses provide information about the extent to which crosslinking, crystallinity or

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Figure 9.20 LTMA analyses of rigid and soft interphases (IP) at filler and elastomer inclusions respectively in a multicomponent polypropylene system.

Tg may be affected by the confinement of the polymer matrix at interfaces. According to L-TMA measurements the inclusion of silica particles or glass fibers in epoxy resin decreases or increases the crosslinking density at the interfacial region depending on the composition of the surfaces. In the case when the reinforcing filler contains adsorbed water, the softening point (and thus the Tg ) in the interfacial region is decreased significantly (due to a reduced density of crosslinking) comparing to the bulk [80]. In contrast, when the adhesion is good the softening transition of the epoxy resin is seen to rise as the fiber is approached, indicating an enhancement of the crosslinking density [81]. The change of interface between acrylnitrile/polybutadiene/styrol terpolymer (ABS) and polyamide (PA) is found to be also strongly adhesion-dependent according to thermal conductivity and L-TMA measurements. Thickness of 50 μm transient region was found around the well-bonded interface, while poor interface caused by insufficient melting of the ABS showed very abrupt transition within a 10 μm thick range. An important issue of micro-thermal analysis is the influence of the topography. False heat flow peaks (increased tip/sample contact area) and troughs (reduced contact area) tend to be produced by valleys and hills, respectively, on the sample surface. Cautious sample preparation and interpretation is required to overcome this problem. The development of nano-thermal analysis by means of extremely fast chip calorimeters was a breakthrough in studying the time-dependent thermal transitions [82, 83]. A differential AC-chip calorimeter is capable of measuring the glass transition in nanometer thin films. Due to the differential setup pJ/K sensitivity is achieved. Heat capacity can be measured for sample masses below one nanogram. The ultra-small sample size makes this calorimeter especially suitable for the frequency-dependent measurement of complex heat capacity in the frequency range from 1 Hz to 1 kHz. The achievable fast (even 1,000,000 Ks-1) heating rates enable the simulation of fast combustion processes and minimizes the changes of metastable samples during measurements. The high cooling rate capability of the system could make a chance for amorphization of semicrystalline polymers, like isotactic polypropylene (iPP) [84] and polyamide-6 (PA6) [85]. Further possibility of micro-thermal analysis belongs to the field of analytical pyrolysis. Localized mass spectrometry of the gases evolved after degradation induced by heated tip has been reported [86, 87].

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9.6

Thermal Gravimetric Analysis and Evolved Gas Analysis

Another large branch of thermal analysis is thermogravimetry. Thermogravimetric Analysis (TGA) measures the amount and rate of change in the weight of a material as a function of temperature or time in controlled atmosphere. TGA is particularly useful for the following measurements: thermal stability, decomposition kinetics, composition, estimated lifetime, oxidative stability, moisture and volatile contents. The temperaturemodulated TGA (MTGA) method is also available for some commercial instruments. MTGA is useful for the kinetic study of the effect of temperature on the decomposition of materials and it can monitor activation energy as a function of conversion, which indicates the mechanism involved. All of these features are relevant for multicomponent polymer systems but mostly the degradation characteristics of polymer blends, composites and fire-retarded systems are analyzed by TGA. The management of polymer wastes (pyrolisis, incineration, recycling, etc.) also initiates several questions that can be answered by thermal gravimetric analysis and its coupled techniques. The heat-induced decomposition or desorption processes can be characterized by their enthalpy and weight loss but the identification of the evolved gaseous products require the analysis of the gas phase (EGA). For this purpose the most frequently-used analytical tools are the mass spectrometer (MS) and Fourier transformed infrared spectrometer (FTIR). In the case of evolved gas of complex composition gas chromatography (GC) is included between the two main units. The off-line ‘Evolved Gas Collection’ technique applies a sorption tube for collecting the gases, the separation and analysis of which occurs in a subsequent step. The on-line techniques, shown in Figure 9.21, identify the components just after the evolution of gases. The most critical part of the coupled systems is the junction (‘interface’). Condensation of the degradation products originating from the thermal decomposition of polymers may occur in this junction. In order to minimize this effect, the rate of carrier gas can be increased which, however, decreases the concentration and residence time of the molecules to be analyzed at the detector. Heating of the connecting tube is the other solution; this, however, increases the chance for secondary decomposition and reaction between the components of the gas. Thus the results are influenced by the length and temperature of junction, flow rate and profile, miscibility with the carrier gas and diffusion coefficient. The applicability of conventional TGA-coupled systems is limited by their maximal heating rate (∼100 ◦ C/min), which is low compared to the real fire scenarios. Therefore these results hardly allow any extrapolation to large-scale fire situations. For increasing the heating rate micro-furnace, high frequency inductance coil (Curie-point pyrolyzer), and rapid voltage control of metals of high resistivity (heated wire pyrolyzer) have been used [88, 89]. More recently, fast-heating TGA equipment uses built-in additional heating lamps or laser sources. The laser method is commonly used for the measurement of the thermal diffusivity and thermal conductivity of solids, powders and liquids. Thermal analytical tools applying laser heat source (laser pyrolyzer) have several advantages over traditional systems. The focused high power ensures rapid enhancement of the temperature of a small area

Figure 9.21 Thermal gravimetric analyzer equipped with controllable nitrogen air mixture inlet; optional rapid pyrolyzer heater and controllable valve of outlet to evolved gas analysis by means of FTIR; MS or GC-MS units.

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followed by rapid cooling, thus avoiding the secondary reactions of the decomposition products in the condensed phase. The decomposition reactions under the rapid heat treatment are kinetically controlled (similarly to the case of combustion) in contrast to the conventional thermodynamically-controlled TGA analyses. The fastest ‘flash pyrolyser’ methods apply extreme high heating rates (103 –104 ◦ C/sec), which can be maintained for 20 s duration. The evolved decomposition products are detected with FTIR. The amount of the analyzed sample is 10–50 mg [90]. A comprehensive study on the applicability of TG-GC-MS for the degradation of polymer systems, including blends and composites, and identification of toxic components was published by Raemaekers, and Bart [91]. A more recent paper reported the advantages and disadvantages of the application of FTIR spectrometer for analysis of the gas evolved during TG analysis [92]. For example, FTIR gas analysis does not need high vacuum and the set-up is much simpler than that of TG-MS. The identification limit (∼10 ppm) in the evolved gases meets most of the requirements; however, the interference of small molecules (H2 O, HCl) in the fine structure of the spectra may cause confusion. The most widely-studied multiphase polymer systems in the recent period are clay nanocomposites. In the case of maleated polypropylene the presence of layered silicate nanoparticles does not change the composition of the degradation products [93], only their evolution is delayed compared to the reference polymer. In contrast, PS nanocomposites release higher amounts of α-methyl styrene, dimer and trimer derivatives [92] than PS alone. A wide range of polymer-clay nanocomposites has been compared (PA6, SAN, ABS, EVA, PS, PMMA, PAN, HIPS, PE, PP) by using TG-GC-MS method [94]. Their degradation mechanism explains their fire-retardancy performance. It was concluded that the delay of evolution of small molecular degradation products, characteristic of the majority of composites, can be explained by the gas barrier effect of the protecting layer formed by the catalytic assistance of the clay nanoparticles. Generation of new degradation pathways (typically random chain scission and radical recombination) is characteristic to the degradation of most of these polymer/clay nanocomposites. If the ratio of intramolecular reactions becomes considerable (e.g., PA6, PS and EVA) it results in improved fire retardancy. The radicals originating from PS and EVA are relatively stabile; consequently, their nanocomposites show a larger cut of the heat release peak than others. The stability of organic compatibilizer molecules depends on the composition of nanoclay. The multi-step degradation of alkyl-ammonium salts on the surface of montmorillonite (of free acidic OH group) leads to the formation of char of twice the amount than on synthetic fluorohectorit (of blocked surface OH groups) [95]. Comparing the pyrolysis products of polymers (PC, PS, ABS, PAN) containing ammonium-polyphosphate (APP) and other phosphorous fire retardants, using MS detector, catalytic chain scission of PC was found due to the effect of ammonia and water, released from APP [96]. Similar analyses made clear that carbon black participates in the hydrogen transfer processes of the degradation of polyolefins (PE, PP and PIB) and influences the ratio of α,ω-dialkenes and n-alkanes [97]. Carbon black, depending on its volatile component content, can promote or retard the decomposition. The laser pyrolysis is the most convenient method for modeling the processes taking place in the ‘dark flame’ region between the surface of the sample and the flame [98]. The primary reactions determining the fire performance of a material are taking place in this zone. The recent development of laser pyrolysis technique and TOF-MS analysis allowed the determination of the kinetic parameters of degradation of various polymers (such as PMMA, PU, PS, PAM, PAN) [99, 100]. In these studies. carbon black was added to the non-charring samples in order to enhance the degradation efficiency of the used lasers of visible light. A disadvantage of all of the methods discussed above is the need for a junction between the site of thermal treatment and gas sensor (and separating) units or use of a sorption tube, where secondary reactions can take place. This shortage has been eliminated by a recently-developed simultaneous laser pyrolysis-FTIR system. In this case, the sample is placed in the FTIR equipment, where the controlled heat treatment is accurately performed by CO2 laser. The real-time detection of the evolved gases promotes better understanding of

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Figure 9.22 laser.

Direct FTIR gas detection and Raman analysis of sample under controlled heat treatment with CO2

the degradation mechanisms [101, 102]. Furthermore, the analysis of the chemical changes in the solid phase is aided by a fiber-optic Raman sensor monitoring the heat-dependent physical and chemical structure (Figure 9.22). The method can be combined with the determination of thermal diffusivity using the wellestablished technique of Laser Flash Analysis (LFA), in which the increase in the sample’s temperature resulting from the absorption from a laser flash is measured (Heat Flow Meter; Thermal Conductivity Tester).

9.7

Conclusions

Thermal analysis of multiphase materials represents an essential part of material science. The discussion of thermal analysis in relation to multiphase polymer systems required the brief description of the well-known fundamentals together with topics that are less exposed in the relevant manuals and reviews. Thermal optical microscopy is, for example, essential for visualization of the thermally-induced polymorphic changes in nucleated and non-nucleated semicrystalline polymers and their blends. The detection of the development of a transcrystalline interlayer around inclusions allows developing polymer systems of well-designed interlayers. The importance of thermal analysis strongly increased with the radical development of dynamic, micro and coupled thermal analytical methods, such as modulated DSC and TG, micro-thermal analysis and new evolved gas analysis methods. By promoting the understanding of these methods, this chapter intends to initiate development strategies for widening the applicability towards even more precise contribution to the design and engineering of multiphase systems. The recent advancements in the experimental techniques and instrumentation can give the answer for the need for analyzing changes of interfacial layers and identifying decomposition products at the very beginning of a degradation process. New perspectives of accelerated heating and cooling techniques suggest a chance for the development of in-line thermal analytical methods for process control. The summary provided in this chapter intends to encourage such activities.

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References 1. J. Paulik, F. Paulik, Complex Thermoanalytical Method for the Simultaneous Recording of T, TG, DTG, DTA, TGT, DTGT, TD and DTD Curves Part I. Development and Characterization of Equipment, Thermochimica Acta 3(1), 13–15 (1971). 2. R.F. Speyer, Thermal Analysis of Materials, (Materials Engineering, Vol. 5, Monograph), Marcel Decker, Inc. New York (1994). 3. B. Wunderlich, Thermal Analysis of Macromolecules, A Personal Review, J. Therm. Anal. Calorim. 89 2, 321–356 (2007). 4. J.M. Haudin, Optical Studies of Polymer Morphology. In Optical Properties of Polymers, (Ed. G.H. Meeten) Chap. 4. Elsevier Appl. Sci., London, New York (1986). 5. J. Karger-Kocsis (Ed.), Polypropylene: Structure, Blends and Composites, Vols 1–3, Chapman & Hall, London (1995). 6. D.R. Norton, A. Keller, The Spherulitic and Lamellar Morphology of Melt-Crystallised Isotactic Polypropylene, Polymer 26 704–716 (1985). 7. J. Varga, Supermolecular Structure of Isotactic Polypropylene. J. Materials Sci. 27 2557–2579 (1992). 8. J. Varga, Crystallization, Melting and Supermolecular Structure of Isotactic Polypropylene. In J. Karger-Kocsis (Ed.) Polypropylene: Structure, Blends and Composites, Vol. I. Structure and Morphology, Chapman and Hall. London (1995). pp. 56–115. 9. B. Lotz, J.C. Wittmann, A.J. Lovinger, Structure and Morphology of Poly(propylenes): a Molecular Analysis, Polymer 37 4979–4992 (1996). 10. J. Varga: β-Modification of Isotactic Polypropylene: Preparation, Structure, Processing, Properties, Application. J. Macromol. Sci., Phys. B 41 1121–1171 (2002). 11. J. Varga, G.W. Ehrenstein, Formation of β-Modification in its Late Stage of Crystallization, Polymer 37 5959–5963 (1996). 12. H. Quan, Z.-M. Li, M.-B. Yang, R. Huang, On Transcrystallinity in Semi-crystalline Polymer Composites. Review article Composite Sci. Technol 65(7-8), 999–1021 (2005). 13. J. Varga, J. Karger-Kocsis, Rules of Supermolecular Structure Formation in Sheared Isotactic Polypropylene Melts J. Polymer Sci. Part B: Polymer Phys. Ed. 34 657–670 (1996). 14. J. Varga, J. Karger-Kocsis, Interfacial Morphologies in Carbon Fibre-Reinforced Polypropylene Composites, Polymer 36 4877–4881 (1995). 15. A.K. Bledzki, J. Gassan, Composites Reinforced with Cellulose-based Fibres. Progress in Polymer Science 24 221–274 (1999). 16. J. Varga, J. Karger-Kocsis, The Difference between Transcrystallization and Shear-Induced Cylindritic Crystallization in Fibre-Reinforced Polypropylene Composites, J. Materials Sci. Letters 13 1069–1071 (1994). 17. F.L. Binsbergen, Heterogeneous Nucleation in the Crystallization of Polyolefins. III. Theory and Mechanism, J. Polym. Sci. Polym. Phys. Ed. 11 117–135 (1973). 18. L.A. Utracki and M.M. Dumoulin, Polypropylene Alloys and Blends with Thermoplastics. In J. Karger-Kocsis (Ed.) Polypropylene: Structure, Blends and Composites, Vol. 2. Chapman and Hall. London (1995), pp. 50–94. 19. Gy. Marosi, Gy. Bertalan, P. Anna, A. Tohl, R. L´agner, I. Balogh, P.F. La Mantia, Thermoanalytical Studies of Nucleating Effects in Polypropylene Composites, I. Liquid Crystal Polymer Containing Polypropylene, J. Thermal Anal. 47 1155–1161 (1996). ´ Liber, G. Belina, Polymer Blends Based on the β-Modification of Polypropylene, 20. A. Menyhard, J. Varga, A. European Polymer J. 41 669–677 (2005). 21. J. Varga, A Menyh´ard: Effect of Solubility and Nucleating Duality of N N -dicyclohexyl-2,6-naphthalenedicarboxamide on the Supermolecular Structure of Isotactic Polypropylene, Macromolecules 40 2422–2431 (2007). 22. A. Thierry, B. Fillon, C. Straupe, B. Lotz, J.C. Wittmann, Polymer Nucleating Agents: Efficiency Scale and Impact of Physical Gelation, Progr. Colloid Sci. 87 28–31 (1992). 23. C. Marco, M.A. Gomez, G. Ellis, J.M. Arribas, Comparative Study of the Nucleation Ability of Third-Generation Sorbitol-Based Nucleating Agents for Isotactic Polypropylene, J. Appl. Polymer. Sci. 84 2440–2450 (2002).

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24. M. Kristiansen, M. Werner, T. Tervoort, P. Smith, M. Blomenhofer, H.W. Schmidt, The Binary System Isotactic Polypropylene/Bis(3,4-dibenzilidene) Sorbitol: Phase Behavior and Optical Properties, Macromolecules 36 5150–5156 (2003). 25. J. Menczel, J. Varga, Influence of Nucleating Agents on Crystallization of Polypropylene I. Talc as a Nucleating Agent, J. Thermal Anal. 28 161–174 (1983). 26. B. Puk´anszky, K. Belina, A. Rockenbauer, F.H.J. Mauer, Effects of Nucleation, Filler Anisotropy and Orientation on the Properties of PP Composites, Composites 25 205–214 (1994). 27. Gy. Marosi, R. L´agner, Gy. Bertalan, P. Anna, A. Tohl, Thermoanalytical Studies of Nucleating Effects in Polypropylene Composites II. Filler and Elastomers Containing Polypropylene J. Thermal Anal. 47 1163–1170 (1996). 28. Z.H. Zhang, Y.J. Tao, Z.G. Yang, K.C. Mai, Preparation and Characteristics of Nano-CaCO3 Supported betaNucleating Agent of Polypropylene, Eur Polymer J. 44 1955–1961 (2008). 29. M.R. Meng, Q. Dou, Effect of Pimelic Acid on the Crystallization, Morphology and Mechanical Properties of Polypropylene/Wollastonite Composites, Mater. Sci. Eng. A 492 177–184 (2008). 30. A. Menyh´ard, J. Varga, The Effect of Compatibilizers on the Crystallisation, Melting and Polymorphic Composition of β-Nucleated Isotactic Polypropylene and Polyamide-6 Blends, European Polymer J. 42 3257–3268 (2006). 31. J.D. Menczel, L. Judonits, R.B. Prime, H.E. Bair, M. Reading, S. Swier, Differential Scanning Calorimetry. In J.D. Menczel and R.B. Prime (Eds), Thermal Analysis of Polymers, Fundamental and Application, John Wiley & Sons Inc., Hoboken, New Jersy (2009). 32. E.A. Turi, Thermal Characterization of Polymeric Materials, Academic Press Inc.: New York (1981). 33. S.Z.D. Cheng, Applications to Polymers and Plastics. In P.K. Gallagher (Ed.), Handbook of Thermal Analysis Calorim, Vol. 3, Elsevier: London (2002). 34. G.W. Ehrenstein, G. Riedel, P. Trawiel, Thermal Analysis of Plastics, Hanser, Munich (2004). 35. B. Wunderlich, Thermal Analysis of Polymeric Materials, Springer, Berlin (2005). 36. V.B.F. Mathot, G. Vanden Poel, T.F.J. Pijpers, Benefits and Potentials of High Performance Differential Scanning Calorimetry (HPer DSC). In M.E. Brown and P.K. Gallagher (Eds), Handbook of Thermal Analysis Calorim; Recent Advances, Techniques and Applications, Volume 5, Ch. 8 (2008), pp. 269–297. 37. M.F.J. Pijpers, V.B.F. Mathot, B. Goderis, R.L. Scherrenberg and E.W. van der Vegte, High-speed Calorimetry for the Study of the Kinetics of (De)vitrification, Crystallization, and Melting of Macromolecules, Macromolecules 35(9), 3601 (2002). 38. G. Buckton, A.A. Adeniyi, M. Saunders, A. Ambarkhane, HyperDSC studies of amorphous polyvinylpyrrolidone in a model wet granulation system, Int. J. Pharm. 312(1–2), 61–65. 39. R.L. Danley, P.A. Caulfield and S.R. Aubuchon, A Rapid-scanning Differential Scanning Calorimeter, American Laboratory 40(1), 9 (2008). 40. B. Fillon, A. Thierry, B. Lotz, J.C. Wittmann, Efficiency Scale for Polymer Nucleating Agents, J. Thermal. Anal. Calorim. 42 721–731 (1994). 41. T. Garoff, V. Virkkunen, P. J¨aa¨ skel¨ainen, T. Vestberg, A Qualitative Model for Polymerisation of Propylene with a MgCl2-Supported TiCl4 Ziegler-Natta Catalyst, Eur. Polym. J. 39 1679–1685 (2003). 42. K. Fukao, S. Foejii, Y. Seruyama N. and S. Tsurutani, Structure Formation and Glass Transition in Oriented Poly(Ethylene Terephtalate). In G. Reiter, G.R. Strobl (Eds) Progress in Understanding of Polymer Crystallization, Springer, Berlin/Heidelberg (2007). 43. A. Romankiewicz, T. Sterzynski, The Lamellar Distribution in Isotactic Polypropylene Modified by Nucleation and Processing, Macromol. Symp. 180 241–256 (2002). 44. J.D. Hoffman, J.J. Weeks, Melting Process and the Equilibrium Melting Temperature of Polychlorotrifluoroethylene, J. Res. Natl. Bur. Stand A-Phys. Chem. 66A 13–28 (1962). 45. B. Monasse, J.M. Haudin, Growth Transition and Morphology Change in Polypropylene, Colloid Polym. Sci. 263 822–831 (1985). 46. J. Xu, S. Srinivas, H. Marand, P. Agarwal, Equilibrium Melting Temperature and Undercooling Dependence of the Spherulitic Growth Rate of Isotactic Polypropylene, Macromolecules 31 8230–8242 (1998). 47. H. Marand, J. Xu, S. Srinivas, Determination of the Equilibrium Melting Temperature of Polymer Crystals: Linear and Nonlinear Hoffmann-Weeks Extrapolations, Macromolecules 31 8219–8229 (1998). 48. S. Fairgrieve, Nucleating Agents, Rapra Review Reports 16 1–132 (2005).

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49. A. Thierry, C. Straupe, J.C. Wittmann, B. Lotz, Organogelators and Polymer Crystallisation, Macromol. Symp. 241 103–110 (2006). 50. D. Alcazar, J. Ruan, A. Thierry, B. Lotz, Structural Matching between the Polymeric Nucleating Agent Isotactic Poly(vinylcyclohexane) and Isotactic Polypropylene, Macromolecules 39 2832–2840 (2006). 51. J. Varga, A. Menyh´ard, Crystallization, Melting and Structure of Polypropylene/Poly(Vinylidene-Fluoride) Blends, J. Therm. Anal. Calorim. 73 735–743 (2003). 52. R.H. Zhang, D. Shi, S.C. Tjong, R.K.Y. Li, Study on the Beta to Alpha Transformation of Polypropylene Crystals in Compatibilized Blend of Polypropylene/polyamide-6, J. Polym. Sci. Pt. B-Polym. Phys. 45 2674–2682 (2007). 53. W. Yang, X.G. Tang, B.H. Xie, Y.P. Zhu, M.B. Yang, M. Hou, Heterogeneous Dispersion of the Compatibilizer in the Injection Molding of Polyamide 6/Polypropylene Blends, J. Appl. Polym. Sci. 113 299–305 (2009). 54. T. Nishi, T.T. Wang, Melting Point Depression and Kinetic Effect of Cooling on Crystallization in Poly(vinylydene fluoride)-Poly (methylmetacrylate) Mixtures, Macromolecules 8 909–915 (1975). 55. P. Xing, L. Dong, Y. An, Z. Feng, M. Avella, E. Martuscelli, Miscibility and Crystallization of Poly(betahydroxybutyrate) and Poly(p-vinylphenol) Blends, Macromolecules 30 2726–2733 (1997). 56. F. Yang, Z.B. Qiu, W.T. Yang, Miscibility and Crystallization of Biodegradable Poly(butylene Succinate-co-butylene Adipate)/poly(vinyl phenol) Blends, Polymer 50 2328–2333 (2009). 57. P. Rathi, T.M. Huang, P. Dayal, T. Kyu, Crystalline-amorphous Interaction in Relation to the Phase Diagrams of Binary Polymer Blends Containing a Crystalline Constituent, J. Phys. Chem. B 112 6460–6466 (2008). 58. H. Liang, F. Xie, W. Wu, F.Q. Guo, B. Chen, J. Xu, Miscibility and Melting Properties of Poly(ethylene 2,6naphthalate)/poly(trimethylene terephthalate) Blends, J. Mater. Sci. 43 2739–2744 (2008) 59. P. Szab´o, B. Puk´anszky, Miscibility of Crystalline and Amorphous Polymers: Poly-ethylene/polyisobutylene Blends, Macromol. Symp. 129 29–42 (1998). 60. A. Menyh´ard, G. Faludi, J. Varga, Beta-Crystallisation Tendency and Structure of Polypropylene Grafted by Maleic Anhydride and its Blends with Isotactic Polypropylene, J. Therm. Anal. Calorim. 93 937–945 (2008). 61. T.G. Fox, Influence of Diluent and of Copolymer Composition on the Glass Temperature of a Polymer System Bull. Am. Phys. Soc. 1 123 (1956). 62. W.N. Kim, C.M. Burns, Blends of Polycarbonate and Poly(Methyl Methacrylate) and the Determination of the Polymer Polymer Interaction Parameter of the 2 Polymers, Macromolecules 20 1876–1882 (1987). 63. E. Fekete, E. F¨oldes, B. Puk´anszky, Effect of Molecular Interactions on the Miscibility and Structure of Polymer Blends, European Polym J. 41 727–736 (2005). 64. P.S. Gill, S.R. Sauerbrunn, M. Reading, Modulated Differential Scanning Calorimetry, J. Therm. Anal. Calorim. 40 931–939 (1993). 65. M. Reading, D. Elliot, V.L. Hill, A New Approach to the Calorimetric Investigation of Physical and Chemical Transitions, J. Therm. Anal. Calorim. 40 949–955 (1993). 66. D.Q.M. Craig, M. Reading, Thermal Analysis of Pharmaceuticals, CRC Press, Taylor & Francis Group, New York (2007). 67. J.D. Menczel, L.H. Judonits, R.B. Prime, H.E. Bair, M. Reading, S. Swier, Modulated Temperature Differential Scanning Calorimetry (MTDSC). In J.D. Menczel, R.B. Prime (Eds), Thermal Analysis of Polymers, Fundamentals and Application, John Wiley & Sons Inc., Hoboken, New Jersey (2009). 68. U. J¨orimann, G. Widmann, R. Riesen, Temperature Modulated DSC (TMDSC) Applications and Limits of Phase Information, Cp Determination and Effect Separation, J. Therm. Anal. Calorim. 56(2), 639–647 (1999). 69. I. Fraga, S. Montserrat, J.M. Hutchinson, Vitrification During the Isothermal Cure of Thermosets, J. Therm. Anal. Calorim 91(3), 687–695 (2008). 70. M. Schubnell, C. Heitz, T. Hutter, S. Sauerbrunn, J.E.K. Schawe, A Multifrequency Temperature-Modulated Technique for DSC, American Laboratory 38 18–21 (2006). 71. J. Bajdik, G. Regdon Jr., T. Marek, I. Er¨os, K. S¨uvegh, K. Pintye-H´odi, The Effect of the Solvent on the Film-forming Parameters of Hydroxypropyl-Cellulose, Int. J. Pharm. 301 192–198 (2005). 72. G.Z. Papageorgiou, A. Docoslis, M. Georgarakis, D. Bikiaris, The Effect of Physical State on the Drug Dissolution Rate. Miscibility Studies of Nimodipine with PVP, J. Therm. Anal. Calorim. 95 903–915 (2009). 73. E. Zuza, J.M. Ugartemendia, A. Lopez, E. Meaurio, A. Lejardi, J.-R. Sarasua, Glass Transition Behavior and Dynamic Fragility in Polylactides containing Mobile and Rigid Amorphous Fractions, Polymer 49 4427–4432 (2008).

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74. S. Argin-Soysal, P. Kofinas, Y. M. Lo, Effect of Complexation Conditions on Xanthan–Chitosan Polyelectrolyte Complex Gels, Food Hydrocolloids 23 202–209 (2009). 75. D. L¨orinczy, F. K¨oncz¨ol, L. Farkas, J. Bel´agyi, C. Schick, Nucleotides Induced Changes in Muscle Fibres Studied by DSC, TMDSC and EPR. J. Therm. Anal. Calorim. 66 633–644 (2001). 76. L. Augsburger, S.W. Hoag, Pharmaceutical Dosage Forms: Tablets Volume 1 Informa Healthcare (2008). 77. N.S. Lawson, R.H. Ion, H.M. Pollock, D.J. Hourston, M. Reading, Characterizing Polymer Surfaces: Nanoidentation, Surface Force Data, Calorimetric Microscopy, Physica Scripta 55 199–205 (1994). 78. A. Hammiche, L. Bozec, M. Conroy, H.M. Pollock, G. Mills, J.M.R. Weaver, D.M. Price, M. Reading, D.J. Hourston, M. Song, J. Vac. Sci. Technol., B 18 1322–1332 (2000). 79. M. Reading, The Use of Modulated Temperature Programs in Tthermal Mmethods, J. Therm. Anal. Calorim. 64 7–14 (2001). 80. G. Van Assche, B. Volckaerts, B. Van Mele, 2000 Presented at 2nd Int. Micro-TA Symp. (Western Kentucky University, USA, 8/9 May 2000). 81. R. Hassler, E. Muhlen, An Introduction to mu TA (TM) and its Application to the Study of Interfaces, Thermochim. Acta. 361 113–120 (2000). 82. M.Y. Efremov, E.A. Olson, M. Zhang, S.L. Lai, F. Schiettekatte, Z.S. Zhang, L.H. Allen, Thin-Film Differential Scanning Nanocalorimetry: Heat Capacity Analysis, Thermochim. Acta 13 412 (2004). 83. A.A. Minakov and C. Schick, Ultrafast Thermal Processing and Nanocalorimetry at Heating and Cooling Rates up to 1 MK/s Rev. Sci. Instrum. 78(7) (2007) 073902. 84. D. Mileva, R. Androsch, E. Zhuravlev, C. Schick, Critical Rate of Cooling for Suppression of Crystallization in Random Copolymers of Propylene with Ethylene and 1-butene, Thermochim. Acta 492 67 (2009). 85. R.T. Tol, A.A. Minakov, S.A. Adamovsky, V.B.F. Mathot and C. Schick, Metastability of polymer crystallites formed at low temperature studied by ultra fast calorimetry: Polyamide 6 confined in sub-micrometer droplets vs. bulk PA6 Polymer 47(6) (2006) 2172. 86. M. Reading, D.M. Price, D.B. Grandy, R.M. Smith, L. Bozec, M. Conroy, A. Hammiche, H. M. Pollock, Microthermal Analysis of Polymers: Current Capabilities and Future Prospects, Macromolecular Symposia 167 45–62 (2001). 87. D.M. Price, M. Reading, R.M. Smith, A. Hammiche, H.M. Pollock, Localised Evolved Gas Analysis by MicroThermal Analysis, J. Thermal Anal. Calorim. 64 309–314 (2001). 88. F.W. Fifield, D. Kealey, Principles and Practice of Analytical Chemistry, 5th ed, Oxford, UK: Blackwell (2000). 89. Z. Parsi, N. Hartog, T. Gorecki, J. Poerschmann, Analytical Pyrolysis as a Tool for the Characterization of Natural Organic Matter – A Comparison of Different Approaches, J. Anal. Appl. Pyrol. 79 9–15 (2007). 90. A Lunghi, L. Gigante, P. Cardillo, V. Stefanoni, G. Pulga., R. Rota, Hazard Assessment of Substances Produced from the Accidental Heating of Chemical Compounds, J. Hazard. Mater. A. 116 11–21 (2004). 91. K.G.H. Raemaekers, J.C.J. Bart, Effect of the Brush Structure on the Degradation Mechanism of Polystyrene-Clay Nanocomposites, Thermochim. Acta. 295 1–58 (1997). 92. K. Chen, M.A. Susner, S. Vyazovkin, Effect of the Brush Structure on the Degradation Mechanism of PolystyreneClay Nanocomposites, Macromol. Rapid Commun. 26 690–695 (2005). 93. A. Tidjani, Polypropylene-Graft-Malefic Anhydride-Nanocomposites: II – Fire Behaviour of Nanocomposites Produced under Nitrogen and in Air, Polym. Degrad. Stabil. 87 43–49 (2005). 94. B.N. Jang, M.C. Costache, C.A. Wilkie, The Relationship between Thermal Degradation Behavior of Polymer and the Fire Retardancy of Polymer/Clay Nanocomposites, Polymer 46 10678–10687 (2005). 95. F. Bellucci, G. Camino, A. Frache, A. Sarra, Catalytic Charring-Volatilization Competition in Organoclay Nanocomposites, Polym. Degrad. Stabil. 92 425–436 (2007). 96. Z. Cz´eg´eny, M. Blazs´o, Effect of Phosphorous Flame Retardants on the Thermal Decomposition of Vinyl Polymers and Copolymers, J. Anal. Appl. Pyrol. 81 218–224 (2007). 97. E. Jakab, M. Omastova, Thermal Decomposition of Polyolefin/Carbon Black Composites, J. Anal. Appl. Pyrol. 74 204–214 (2005). 98. D. Price, J.R. Ebdon, T.R. Hull, G.J. Milnes, B.J. Hunt, Studies of Chemical Behaviour in Different Regions of Polymer Combustion and the Influence of Flame Retardants Thereon, Fire Polym. 24 307–332 (2001).

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99. D. Price, F. Gao, G.J. Milnes, B. Eling, C.I. Lindsay, T.P. McGrail, Laser Pyrolysis/Time-of-Flight Mass Spectrometry Studies Pertinent to the Behaviour of Flame-Retarded Polymers in Real Fire Situations, Polym. Degrad. Stabil. 64 403–410 (1999). 100. J.R. Ebdon, D. Price, B.J. Hunt, J.P. Gao, F. Milnes, G.J. Cunliffe, L.K. Flame Retardance in some Polystyrenes and Poly(methyl methacrylate)s with Covalently Bound Phosphorus-containing Groups: Initial Screening Experiments and some Laser Pyrolysis Mechanistic Studies, Polym. Degrad. Stabil. 69 267–277 (2000). 101. B. Marosfoi, G. Marosi, A. Szabo, B. Vajna, A. Szep, Laser Pyrolysis Micro-Spectroscopy for Modelling Fireinduced Degradation of Ethylene-Vinyl Acetate Systems, Polym. Degrad. Stabil. 92 2231–2238 (2007). 102. B. Bodzay, B.B. Marosfoi, T. Igricza, K. Bocz, G. Marosi, Polymer degradation Studies using Laser Pyrolysis-FTIR Microanalysis, J. Anal. Appl. Pyrolysis 85 1–2, 313–320 (2009).

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10 Thermophysical Properties of Multiphase Polymer Systems Abderrahim Boudenne, Laurent Ibos and Yves Candau Universit´e Paris-Est Cr´eteil, Certes 61 Av. du G´en´eral de Gaulle 94010 Cr´eteil Cedex, France

10.1 Introduction Polymers are materials with low values of thermal conductivity, which roughly varies from 0.15 W.m−1 .K−1 for amorphous polymers to the 0.5 W.m−1 .K−1 for highly crystalline polymers [1]. In various industrial applications there are many reasons to increase thermal properties, and particularly thermal conductivity of polymer-based materials. The applications are associated in different fields of engineering (microelectronics, aeronautics and space, transport . . .). One way to improve the thermal behavior of these materials is to combine polymeric matrices like thermoplastic polymers (polypropylene, polyamide, etc.), thermosetting resins (epoxy, polyurethane) or rubbers with conductive fillers. Metallic or inorganic materials and graphite in several configurations (particles, fibers, nanotubes) are commonly used as thermally-conductive fillers [2–5]. The association of polymer and fillers allows obtaining composites whose properties are intermediate between one of the two components. So, the interest for these materials arises from the fact that it is possible to develop new materials with properties adapted to specific applications. Moreover, in some cases, composite materials allow the physical properties of each component used in the manufacturing process to be combined. In the different types of properties of composites, thermophysical properties such as effective thermal conductivity, thermal diffusivity, thermal effusivity and specific heat are of paramount importance for the scientists and engineers for specific use of a particular composite for a specific purpose [6–8]. Information on the thermophysical properties of polymer blends or polymer composite materials is also necessary for determining optimum conditions during processing of materials, as well as for analyzing heat transport in materials during practical applications [9]. It was shown long ago that spatial distribution of the dispersed metallic fillers influences thermophysical behavior of the composites [10, 11]. When filler distribution results

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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in formation of a segregated structure, the thermophysical properties get enhanced as compared to a random distribution of the filler in the polymeric matrix. The improvement of the thermal properties of polymer composite material is strongly dependent on several factors, such as orientation and distribution of fillers in the polymeric matrix, the manufacturing process, the surface contacts and adhesion between two components [12]. The problem posed here is how to achieve a reliable prediction of composite properties based on the knowledge of the properties of each isolated component and the concentration of each component in the mixture. Only a few publications have been devoted to the correlation between their structure and their thermal properties [13]. Two-phase systems have been proposed to predict thermal properties with respect to filler concentration, but generally these models are only consistent for low concentrations due to influent parameters that have not been taken into account in the models [14].

10.2 Thermophysical Properties: Short Definitions Heat transfer through a material may involve three modes of transfer: conduction, convection and radiation. By definition, heat transfer by conduction is achieved without displacement of matter but is only due to the presence of temperature gradients between different regions of the material. In opaque solid materials, conduction is the only heat transfer mode to consider. Convection heat transfer has to be considered in materials in liquid state or in porous media. This heat transfer mode involves the displacement of matter from one point to another, so it can act only in fluid-like gases or liquids. The third heat transfer mode is related to radiation phenomena; radiative heat exchanges have to be considered in porous media or in semi-transparent materials. In most cases and applications, composites materials or polymer blends can be considered as opaque solid materials. Thus conductive heat transfer is the main heat transfer to consider. Moreover, when radiative or convective heat transfers are also present (porous materials, for instance) effective properties can be considered. In that case, these properties take into account the different heat transfer modes inside the considered material. In thermal science the term ‘thermal properties’ may indicate a wide range of definitions, mechanisms and phenomena. However, in this chapter, we will focus our interest on the thermophysical properties, since this term includes some specific thermal properties which are: thermal conductivity, thermal diffusivity, thermal effusivity, and specific heat. These properties are always important and often critical in both processing stages and product uses of materials. In the case of composite or multi-component systems the effective (eff ) term is added to the thermophysical property defined simply in terms of the averages of various quantities over the system. Thermal conductivity (k) can be defined as the material property which indicates the quantity of heat transferred during one second through a section of 1 m2 of a material with a thickness of 1 m when the temperature difference between two opposite sides is equal to 1 K (Figure 10.1). This quantity is always a

T

T+1

1s

k

1m

Figure 10.1

Schema of description of thermal conductivity.

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positive one and is expressed in SI units as W.m−1 .K−1 and in cgs units as cal/(cm.s.◦ C). In insulating materials thermal conduction is mainly induced by phonon vibrations. On the contrary, in electrical conducting materials thermal conduction is mainly due to electron displacements. Specific heat (Cp) represents the heat amount required to raise the temperature by one degree of a mass unity of a material. More heat energy is required to increase the temperature of a substance with high specific heat capacity than when using a low specific heat capacity material. Dividing heat capacity by the body’s mass yields a specific heat capacity, which is no longer dependent on the amount of material but on the type of material and on the temperature. Two specific heats can be defined: one at constant volume (Cv) and the other one at constant pressure (Cp). Most engineering applications use specific heat at constant pressure rather than at constant volume. The Cp is expressed in SI units as J.kg−1 .K−1 and in cgs units as cal/(g. ◦ C). Thermal diffusivity (α) describes the way heat flows through a material. In fact, thermal diffusivity is used since it is generally easier and faster to measure than thermal conductivity. It is also sometimes defined as the rate of change of temperature in a transient heat transfer process. This property is expressed in SI units as m2 /s and in cgs units as cm2 /s. This thermal property is somewhat neglected and unknown in polymer and composite sciences as opposed to both thermal conductivity and specific heat (Cp). Thermal diffusivity can be computed from thermal conductivity and specific heat values by: α = k/ (ρCp), where ρ is the material density. Thermal effusivity is probably the least considered thermophysical property in polymer and composites science. By definition, thermal effusivity is the property that fixes the interfacial temperature when a contact is formed between two materials objects initially at different temperatures.√Effusivity combines thermal conductivity, density and heat capacity into one value and is defined by e = kρCp = √kα with dimension W.s1/2 .m−2 .K−1 . Actually, thermal effusivity is a measure of the sample’s thermal impedance or its ability to exchange heat with the environment. Thermal effusivity is a relevant thermophysical parameter for modeling heat exchanges in thermal treatment processes (quenching for instance).

10.3 Measurement Techniques The thermophysical characterization of materials is an important field of metrology which concerns the determination of thermophysical properties related. In the general case, the four thermophysical parameters are difficult to identify using a single experiment. A large number of experimental methods have been developed for determining the thermophysical properties of materials. In conventional techniques, the measurement of each thermophysical property is obtained using specific devices and methods [15, 16]. The heat transport equation is used in order to find thermophysical parameters by the measurement of heat fluxes or temperature gradients [17]. In this section, we have chosen to classify the measurement methods in two categories: measurement methods that allow measuring a single thermophysical property and measurement methods that allow measuring several parameters at once. 10.3.1 10.3.1.1

Methods for the Measurement of One Property Thermal Conductivity

Thermal conductivity is usually measured in steady state conditions. Nevertheless, transient or periodic methods can also be applied for the measurement of this parameter. Whatever the measurement technique, the determination of this parameter is based on the following methodology. One side of the sample (sometimes the whole sample) is subjected to a constant or variable heat flux during duration depending on the material

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under study and on its sample shape and dimensions. The study of the evolution of temperature as a function of time allows the calculation of the thermal conductivity. The choice of a method for measuring the thermal conductivity of the samples may depend on the material (granular, semi-infinite, . . .), shape (2D, 3D, . . .), tolerance (or accuracy) of certain parameters. Several methods have been developed for the measurement of the thermal conductivity. However, in this section we will focus the presentation on the most common methods such as the Guarded Hot-plate method, Hot wire method, and 3ω method. 1. Guarded Hot-plate method: This static method is based on the heating of one side of the sample with a unidirectional constant and known heat flow, and on the measurement of the temperature difference between sample faces. The technique has subsequently become very well established and is documented in the written standard ISO8302 of 1991. It is now unarguably recognized as the most accurate technique for determining the thermal conductivity of insulations, having an uncertainty of about 1.5% over a limited temperature range near ambient [18]. 2. The Hot wire method is one of the best known and most used measurement methods of the thermal conductivity. The method is documented in the written standard ISO8894. Initially, this method has been developed for the measurement of conductivity of liquids, and was later extended to the characterization of solids and porous media. The linear heat source (hot wire) is embedded in the tested material. If the heat source is assumed to have a constant and uniform dissipation along the wire embedded in the sample, the thermal conductivity can be derived directly from the resulting change in the wire temperature over a known time interval. One of the interests of this method is that the wire temperature can be estimated by the measurement of its electrical resistivity. So the heating source is also used as a temperature sensor. 3. The 3ω method was reported for the first time for the measurement of the thermal conductivity of solids by Cahill [19]. A micro-fabricated metal line is deposited onto the specimen to act as a heater/thermometer. When an alternating current (AC) voltage signal is used to excite the heater at a frequency ω, the periodic heating generates oscillations in the electrical resistance of the metal line at a frequency of 2ω. In turn, this leads to a third harmonic (3ω) in the voltage signal, which is used to measure the magnitude of the temperature oscillations. The frequency dependence of the oscillation amplitude and phase can be analyzed to obtain the thermal conductivity of the specimen. Thermal Conductivity for Orthotropic Composites For orthotropic composite and knowing the principal axes, the components of the thermal conductivity tensor can be found with the simplest method. For 3D problems, it consists of measuring the thermal conductivity along each direction with three samples using 1D periodic or steady state measurement techniques [20]. For the in-plane directions, two samples are usually realized by cutting the initial sample into several equal strips, rotating them by 90◦ and reintegrating them to obtain new samples [20]. Otherwise, simultaneous measurement techniques are available. One can distinguish:

r r

numerical and inverse methods with thermocouples embedded in samples and with plane or cylindrical heat sources [21–23] analytical heat transfer solutions with temperature field provided by infrared detector or camera and with laser heating [24, 25].

10.3.1.2

Thermal Diffusivity

The thermal diffusivity of the materials is more often measured using the flash method (see Figure 10.2). In the original technique a flash of radiant energy is deposited over the front surface of a homogeneous sample

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Instantaneous pulse Front face Initial temperature T0 Thickness L Increase of temperature T0 + ΔT(t)

Rear face

Figure 10.2

Schematic of the flash method.

and the temperature increase T(t) on the rear face of the sample is recorded as a function of time. The main interests of this method are, on the one hand, its short measurement duration and, on the other, the fact that this method does not use any contact with the sample if the sample rear face temperature is measured with non-contact sensors. The heating source used initially was a flash lamp (discharge source). Recently, improvements to the method have led some authors to use laser sources. This last improvement allows obtaining values of thermal diffusivity depending on the direction. This can be useful for the characterization of non-isotropic samples as, for instance, polymers reinforced with woven fibers. Typical shapes of the temperature increase on the rear side of the sample are presented Figure 10.3. If no heat losses are involved, the temperature of the rear face will rise to a maximum and remain at that level indefinitely (curve 1). However, with increasing heat losses, the temperature on the rear face decreases after reaching a maximal value (curves 2 and 3). The original method proposed by Parker et al. [26] assumes that heat losses are negligible. Hence, the thermal diffusivity can be determined using: α = 0.139 ×

L2 t1/2

(10.1)

where parameter L is the thickness of the sample and t1 /2 is the time necessary to reach half of the maximal temperature increase. Since this method assumes ideal conditions of adiabatic sample and instantaneous pulse heating, it is somewhat limited in applicability. To make it more suitable to experimental conditions, other methods have

ΔT

(1)

TM h TM/2

(2) (3)

0

t1/2

t

Figure 10.3 Rear face temperature increase during a flash method experiment for various experimental conditions: (1) no heat exchanges; (2) and (3) increasing heat exchanges.

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been introduced over the years, which account for heat losses, finite pulse duration, nonuniform pulse heating and composite (non-homogeneous) structures [27]. However, the mathematical analysis in this case is more complex and requires the use of inverse methods for the identification of thermal diffusivity [28, 29]. Specific Heat Capacity The determination of heat capacity of a sample of mass m requires the measurement of the heat amount dQ necessary to increase the temperature of the sample of dT: Cp =

1 dQ m dT

(10.2)

and keeping constant the composition of the material and the pressure. Classically, the measurements of specific heat are performed with an adiabatic calorimeter [30]. Even today, adiabatic calorimetry is the most accurate measurement method for low temperatures. In an adiabatic calorimeter an attempt is made to follow step-wise temperature changes of an internally heated calorimeter in a well-controlled, adiabatic enclosure. Due to heat losses caused by deviations from adiabatic conditions, corrections must be made to the heat added to the calorimeter, dQ. Similarly, the measured increase of temperature, dT, must be corrected for the temperature drifts of calorimeter, where: Cp =

Q corrected − C  Tcorrected mTcorrected

(10.3)

Cp is calculated by subtracting the heat capacity of the empty calorimeter, its ‘water value’ C , and dividing by the mass of the sample, m [30]. Modern control and measurement technology developed over the last 30 years allows miniaturized calorimeters to measure down to milligram quantities in differential scanning mode (DSC) [30, 31]. Differential scanning calorimetry is a modification of differential thermal analysis. In the DSC technique, the device records the power difference necessary to impose a similar given temperature program to the tested sample and to a reference sample. These two samples are placed in separate furnaces (Figure 10.4). A calibration of the device using a material sample of known specific heat is required. The use of samples with low mass (a few milligrams) allows obtaining a rapid characterization in a broad temperature domain. The DSC method is highly suitable to polymer materials and can be used to measure not only specific heat capacity but

Reference

Sample

Temperature control

Computer Power

Cooling system

Figure 10.4

Schematic diagram of differential scanning calorimetry device.

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Reference

Flash lamp

Flash lamp IR Camera

Figure 10.5

Schematic diagram of the flash method.

also to observe transitional phenomena (melting, crystallization, phase transitions, polymerization . . .) and to measure some related parameters such as melting, crystallization or reaction enthalpies, rates of melting, crystallization or reaction . . . [32]. More information about this technique and its application to polymers, blends or composites is provided in Chapter 9 of this book. Thermal Effusivity The measurement of the thermal effusivity requires the perturbation of the thermal balance of the material and the recording of the temperature of the heated sample. Two types of measurement methods can be found in the literature. Some techniques are based on the use of a hot source placed in contact with the sample surface: the most commonly-used measuring method in this case is the Touchau method which was modified to obtain more accurate measurements [33]. The second way is based on non-contact measurement methods. In this case, the sample is submitted to a radiative excitation (laser, halogen lamp, . . .) and the observation of the evolution of the heated surface temperature allows estimating sample thermal effusivity. The sample surface temperature can be measured by temperature sensors or by using an infrared camera. Several methods are referenced in the literature and the most common one uses flash lamps (see Figure 10.5). Measurements can also be carried out by comparison with a reference material, thus avoiding exact knowledge of the absorbed power density and calibration of the detector. 10.3.2

Methods for the Simultaneous Measurement of Several Parameters

In the last 20 years, several techniques have been developed for the simultaneous measurement of thermophysical properties of materials [34–36]. Photoacoustic and photopyroelectric techniques are two measurement methods which use a laser source as a periodic heating device and allow simultaneous estimation of the thermophysical parameters [37]. In the photoacoustic method, heat propagation is detected acoustically by a microphone and thermal conductivity and diffusivity are simultaneously estimated. Some restrictions were made concerning the accuracy of the photoacoustic method compared to the photopyroelectric technique [38]. These authors showed that the latter method gives more accurate information concerning both amplitude and phase than the former. In the photopyroelectric technique a pyroelectric transducer detects temperature variations. Both thermal conductivity and specific heat are estimated simultaneously. The knowledge of density is then necessary to find the thermal diffusivity. The transient plane source (TPS) technique is a well-known technique for the measurement of thermal parameters [39]. The hot disk device (see Figure 10.6) applies the TPS technique and can be used over quite a broad range of temperatures (from 30 to 1000 K) and allows the measurement of the thermal conductivity of insulating or conducting materials (between 0.01 and 500 W.m−1 .K−1 ). This technique is an extension of

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Radial

Sensor

Figure 10.6

Sample

Hot disk measurements.

the transient hot strip (THS) method developed by Gustafsson et al. in 1979, which is based on a line-heat source method [40]. Other methods belonging to the class of transient heat flow have been reported for the simultaneous determination of thermal effusivity and diffusivity of materials such as the extended dynamic plane source (DTP) technique proposed for materials of low thermal conductivity. This method is derived from the hot symmetrical plane source by the fact that a material of high effusivity acting as a heat sink is added to both sides of the device [41]. The transient hot wire (THW) and hot-wire parallel technique [42–44] are also largely used to measure thermal conductivity and diffusivity of materials, particularly of polymers (Figure 10.7). Recently, Boudenne et al. [45, 46, 48] have proposed a new periodical method (DICO method) to estimate simultaneously thermal conductivity and diffusivity of polymer and composite materials. This measurement method does not require the knowledge of the sample heat capacity or density and is based on the use of small temperature modulations at different frequencies inside the sample. An extension of this method was proposed by Agoudjil et al. [48] to estimate also sample emissivity: in that case, the method requires the use of an infrared camera as temperature sensor.

10.4 Thermophysical Properties of Polymers and Composite Systems 10.4.1

Neat Polymers (Unfilled Systems)

The thermal conductivity and diffusivity of ceramics and electrically-conducting material such as metals are higher than those of polymers. In fact, pure polymers have low thermal conductivity, ranging from 0.1 to 0.6 W.m−1 .K−1 . Generally, crystalline polymers have higher conductivities and diffusivities than amorphous polymers, reflecting the increased order and higher density, which provides more optimal pathways for thermal transport [32]. Polymer foams have lower thermal properties than classical polymers relating to their lower density and to their highly porous structure. Most foamed polymers (expanded or extruded polystyrene,

Hot Wire Sample Thermocouple

Figure 10.7

Parallel hot-wire technique

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polyurethane foam . . .) have thermal conductivity values in the order of 2−5 × 10−2 W.m−1 .K−1 which is about 10 times less than the same neat polymers in compacted state. The effect of temperature on thermophysical properties for a given polymeric system is very important for processes and applications. The temperature dependence of thermal conductivity and diffusivity of polymers has been studied from low temperature to melting point [49–52]. The systematic studies of this dependence are shown in Table 10.1 [32, 52]. Generally, for amorphous polymers, thermal conductivity increases gradually in the glassy region and decreases slowly or remains constant in the rubber region. For crystalline polymers, thermal conductivity decreases steadily with the increase in temperature below the melting point. In the liquid state, thermal conductivity of polymers increases with temperature. This behavior is illustrated schematically in Figure 10.8. Usually, the thermal conductivity of polymers is highly dependent on polymer chain segment orientation. Crystalline polymers have highly ordered chain segments and therefore have higher ability for heat transport than amorphous polymer. Contrarily to the thermal conductivity behavior, the thermal diffusivity of amorphous polymers is seen to decrease slowly in the glassy region, remains constant in the rubbery region and decreases again slowly in the melt-flow region. For crystalline polymers, the thermal diffusivity has a similar behavior to the thermal conductivity in the solid state, as illustrated in Figure 10.9. In the liquid state, a rapid increase of thermal diffusivity is observed followed by a slow decrease. Generally, the different values of specific heat capacity of polymers reported in the literature show at the maximum a factor of variation three [32]. The difference is associated with the different molecular composition of the polymer. The molecular architecture (crystalline or amorphous polymer) has a small influence on specific heat. The temperature dependence of specific heat capacity for a crystalline polymer is reported in Figure 10.10. At the glass transition temperature a small step is observed. This same behavior is observed in amorphous polymers. The melting of crystalline regions of semi-crystalline polymers induces the presence of a relatively large peak; the area of this peak allows computing the melting enthalpy of the polymer. 10.4.2

Thermophysical Behavior of Composites

Thermal conductivity is one of the thermophysical quantities most used in the field of thermal science to model heat transfers in composite materials. For over a century, many relationships have been proposed to predict the effective thermal conductivity of composite materials. ‘Effective’ means that the property is a macroscopic property characteristic of the composite considered as a statistically homogeneous material [57]. Many theoretical, semi-empirical, empirical and numerical models have been developed to predict the value of this effective conductivity based on the characteristics of each component of the composite and sometimes on the dispersion and shape of the inclusions. Bibliographical summaries of these models have been proposed by Mottram and Taylor [57], Gorringe and Churchill [58], Godbee and Ziegler [59], and Cheng and Vachon [60]. Obviously, all effective conductivities obtained with these relationships are necessarily intermediate between the thermal conductivity values of components constituting the composite. It is important to notice that development of new models (theoretical, numerical or semi empirical) need experimental data bases. Table 10.2 summarizes data on the thermal conductivity, thermal diffusivity and specific heat of some polymer composites [32, 61–66]. 10.4.2.1

Theoretical Prediction Models

Theoretical models are established on the basis of physical modeling of transport phenomena in a heterogeneous material. The result of this modeling is a mathematical relationship involving at least the respective volume fractions and the properties of each component of the material. Some models take into account the shape, dispersion and orientation of the fillers within the polymeric matrix.

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Table 10.1 Thermal conductivity of some polymers, polymers blends, polymer composites materials [32, 52–56].

Material

Temperature (K)

Thermal conductivity W.m−1 .K−1

Thermal diffusivity mm2 /s

Specific heat capacity kJ.kg−1 .K−1

Polyamides Polyauryltactam (nylon-12) Polycaprolactam (nylon-6) Moldings Crystalline Amorphous Melt Poly(hexamethylene adipamide) (nylon-6,6) Moldings Crystalline Amorphous Melt Poly(hexamethylene dodecanediamide) (nylon-6,12) Poly(hexamethylene sebacamide) (nylon-6,10) Polyundecanolactam (nylon-11)

0.19–0.25 293 303 303 523

0.24 0.43 0.36 0.24

239 303 303 523

0.24 0.43 0.36 0.15

1.675

0.101

1.67

0.22

293

0.22

1.675

0.23

Polycarbonates, polyesters, polyethers and polyketones Polyacetal Polyaryletherketone Poly(butylene terephthalate) (PBT)

293 293

Polycarbonate (Bisphenol A) Poly(dially carbonate) Poly(2,6-dimethyl-1,4-pheneylene ether) (PPO) Polyester Cast, rigid Chlorinated Polyetheresteramide

293

Polyetheretherketone (PEEK) Poly(ethylene terephthalate) (PET) Poly(oxymethylene) Poly(phenylene oxide) Molding grade

300

0.30 0.23 0.30 0.29 0.16 0.20 0.21

1.6134 0.14

1.255 1.2459

0.12

293

0.17 0.33 0.24–0.34 0.20–0.26 0.25 0.15

0.14

1.425

293

0.292

0.23

1.47

303 353

0.23

1.045

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Table 10.1 (Continued)

Material

Temperature (K)

Thermal conductivity W.m−1 .K−1

293 293 293 293 300–500

0.19 0.84 1.10 0.16 0.19–0.34

293 311–460

0.29 0.146–0.24

Thermal diffusivity mm2 /s

Specific heat capacity kJ.kg−1 .K−1

Epoxides Epoxy resin Casting grade silica filled aluminium filled Molding grade Temperature dependence

1.045

Halogenated olefin polymers Polychlorotrifluoroethylene Poly(ethylene-tetrafluoroethylene) Copolymer Polytetrafluoroethylene

Poly(tetrafluoroethylenehexafluoroethylene) Copolymer(Teflon EEP) Poly(vinyl chloride) (PVC) Rigid Flexible Chlorinated Temperature dependence

Poly(vinylidene chloride) (PVC2) Poly(vinylidene fluoride) (PVF2)

0.85945

0.238 293 298 345

0.25 0.25 0.34

0.095 1.045

0.202 293 293 293 103 273 373 250 293 250 293 298–433

0.21 0.17 0.14 0.129 0.158 0.165 0.13 0.13 0.17–0.19

0.116 0.07

1.000 1.675

1.569 0.7115 1.340 0.7856 1.380

Hydrocarbon polymers Polybulene Polybutadiene Extrusion grade Poly(Butadiene-styrene) copolymer (SBR) 23.5% Styrene content Pure gum vulcanizate Carbon black vulcanizate

0.22 293

0.22

0.190–0.25 0.300 (Continued)

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Table 10.1 (Continued)

Material Polychloroprene (Neoprene) Unvulcanized Pure gun vulcanizate Carbon black vulcanizate Polycyclooctene 80% trans content Poly(1,3-cyclopentylenevinylene) [poly(2-norbornene)] Poly(ethylene) Low density Medium density High density Poly(ethylene-propylene) copolymer Polyisobutylene Polyisoprene (natural rubber) Unvulcanized Pure gun vulcanizate Carbon black vulcanizate Poly(4-methyl-1-pentene) Polypropylene Polystyrene

Poly(p-xylylene) (PPX)

Temperature (K) 293 293

Thermal conductivity W.m−1 .K−1

Thermal diffusivity mm2 /s

Specific heat capacity kJ.kg−1 .K−1 2.175

0.19 0.192 0.210 0.27 0.29

293 293

293

293 293 273 373 473 573 673 300

0.33 0.42 0.52 0.355 0.13 0.13 0.15 0.28 0.167 0.12 0.105 0.128 0.13 0.14 0.160 0.12

0.14 0.23

2.300 2.300 2.075

1.885

0.12 0.13

2.350 1.2230 1.9322 2.4417 1.3022

Polyamides Polyetherimide Polyimide Thermoplastic Thermoset

0.07 293

0.11 0.23–0.50

Phenolic resins Poly(phenol-formaldehyde) resin Casting grade Molding grade Poly(phenol-furfural) resin Molding grade

0.192

293

0.15 0.25

1.465 1.465

0.25

1.460

Polysaccharides Cellulose Cotton Rayon

0.071 0.054–0.07

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Table 10.1 (Continued)

Material Sulphite pulp, wet Sulphite pulp, dry Laminated Kraft Paper Alkali cellulose Different papers Cellulose acetate Cellulose acetate butyrate Cellulose nitriate Cellulose propionate Ethylcellulose Polysilloxanes Poly(dimethylsiloxane) (PDMS)

Poly(methylphenylsiloxane) 9.5% Pheny, d = 1,110 kg/m3 48% Pheny, d = 1,070 kg/m3 62% Pheny, d = 1,110 kg/m3

Temperature (K)

303–333 293 293

Thermal conductivity W.m−1 .K−1

Thermal diffusivity mm2 /s

Specific heat capacity kJ.kg−1 .K−1

0.8 0.067 0.13 0.046–0.067 0.029–0.17 0.20 0.33 0.29 0.20 0.21

230 290 340 410

0.25 0.20 0.20 0.17

273 323 373 273 323 373 273 323 373

0.158 0.150 0.144 0.143 0.136 0.127 0.141 0.137 0.132 1.295

293 240–310

0.18 0.18 0.29 0.288 0.18 0.26

293 293

0.21 0.31

1.800 1.775

293

0.26

1.297

1.591 1.657

Polysulfide and polysulfones Polyarylsulfone Polyethersulfone Poly(phenylene sulfide) Poly(phenylene sulfone) Udel polysulfone Polyurethanes Polyurethane Casting resin Elastomer Vinyl polymers Polyacrylonitrile Poly(acrylonitrile-butadiene) copolymer (NBR)

(Continued)

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Table 10.1 (Continued)

Material 35% Acrylonitrile Poly(acrylonitrile-butadiene-styrene) copolymer (ABS) Poly(acrylonitrile-styrene) copolymer Poly(i-butyl methacrylate) At 0.82 atm Poly(n-butyl methacrylate) At 0.82 atm Poly(butyl methacrylate-triethylene glycol dimethacrylate) copolymer Poly(chloroethylene-vinyl acetate) copolymer Poly(dially phthalate) Poly(ethyl acrylate)

Poly(ethyl methacrylate)m At 0.82 atm Poly(ethylene vinyl acetate) Poly(methyl methacrylate) Poly(methyl methacrylate-acrylonitrile) copolymer Polymer(methyl methacrylate-styrene) copolymer Poly(vinyl acetate) (PVAc) Poly(vinyl acetate-vinyl chloride) copolymer Poly(vinyl alcohol) (PVA) Poly(N-vinyl carbozole) Poly(vinyl fluoride) (PVF)

Poly(vinyl formal) Molding grade

Temperature (K) 413 333

293 300 300

293 325 375 300 422.1 500 300 273 293

Thermal conductivity W.m−1 .K−1

Thermal diffusivity mm2 /s

Specific heat capacity kJ.kg−1 .K−1

0.184 0.251 0.33

0.17

1.47

0.18

0.081

1.38 1.5710

0.13

1.8524

0.45 0.15

0.134 0.146 0.218 0.21 0.213 0.230 0.213 0.175 0.34 0.21 0.18

1.7867 2.2189 1.4666

0.13

1.3755

0.12–0.21

300

0.159 0.167

1.183

300 293 443 243 300 333

0.2 0.126 0.168 0.14

1.546

293

0.27

1.301 0.17

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Thermophysical Properties of Multiphase Polymer Systems Thermal conductivity

(B): partially crystalline polymer

(A): amorphous polymer

Temperature

Figure 10.8

Thermal conductivity vs temperature based on [32].

Thermal diffusivity (B): partially crystalline polymer

(A): amorphous polymer

Temperature

Figure 10.9

Thermal diffusivity vs temperature based on [32].

Specific heat capacity

Tg

Figure 10.10

Tm

Temperature

Specific heat vs temperature based on [32].

401

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Table 10.2 Thermal conductivity, thermal diffusivity and specific heat of some polymer composites [32, 61, 62–65, 66]. Material

Thermal conductivity W.m−1 .K−1

Polypropylene + 5% Cu40 μm 10% Cu40 μm 15% Cu40 μm 20% Cu40 μm 30% Cu40 μm 40% Cu40 μm

0.24 0.29 0.38 0.58 0.94 1.65 2.14

1.66 2.21 2.60 3.61 4.80 7.59 9.82

1731 997 945 820 796 673 600

Polypropylene + 5% Cu420 μm 10% Cu420 μm 15% Cu420 μm 20% Cu420 μm 30% Cu420 μm 40% Cu420 μm

0.29 0.38 0.58 0.94 1.65 2.14

2.21 2.60 3.61 4.80 7.59 9.82

1231 994 865 680 688 678

Polypropylene + 10% Al25 μm 20% Al25 μm 30% Al25 μm 40% Al25 μm 50% Al25 μm 60% Al25 μm

0.25 0.40 0.58 0.98 1.95 2.68

1.70 1.57 3.39 5.33 9.02 11.78

1528 1267 1255 1183 1263 1188

Polypropylene + 10% Al125 μm 20% Al125 μm 30% Al125 μm 40% Al125 μm 50% Al125 μm 60% Al125 μm

0.26 0.51 0.72 1.55 2.59 4.12

2.48 4.12 4.91 9.60 1.32 1.97

1154 1015 1098 1031 1127 1086

2.91

1676

2.59 2.97 3.21 4.07 5.17 6.58

1696 1511 1396 1372 1229 1134

0.12

1.08

1235

0.16 0.21 0.30 0.38 0.46 0.50

1.24 1.34 1.62 1.99 2.58 2.77

1552 1623 1836 1726 1511 1432

HDPE + 0.45 Silver coated Polyamide particles = PA-Ag 10% PA-Ag 0.42 20% PA-Ag 0.49 30% PA-Ag 0.55 40% PA-Ag 0.76 50% PA-Ag 0.97 60% PA-Ag 1.37 EVA (18%Va) + Expensed Graphite = EG 3% EG 7% EG 10% EG 15% EG 22% EG 30% EG

Thermal diffusivity (m2 /s)· 10−7

Specific heat capacity J.kg−1 .K−1

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Table 10.2 (Continued) Thermal conductivity W.m−1 .K−1

Thermal diffusivity (m2 /s)· 10−7

EVA (18%Va) + Unexpended Graphite = UG 2% UG 7% UG 10% UG 16% UG 23% UG 32% UG

0.12

1.08

109

0.18 0.28 0.25 0.28 0.27 0.43

1.09 1.44 1.56 1.94 2.51 3.24

256 196 116 183 149 202

EVA (18%Va) + Glass sphere = G 3% G36 μm 8% G36 μm 15% G36 μm 29% G36 μm 36% G36 μm 45% G36 μm

0.27

1.17

2538

0.27 0.29 0.33 0.37 0.45 0.48

1.18 1.26 1.54 1.97 2.64 2.38

2369 2176 1961 1343 1307 1198

EVA (18%Va) + 0.27 Silver coated glass sphere = GAg 0.28 4% GAg14 μm 0.33 10% GAg14 μm 0.36 15% GAg14 μm 0.42 21% GAg14 μm 0.65 31% GAg14 μm 0.85 41% GAg14 μm 1.09 43% GAg14 μm 1.38 50% GAg14 μm

1.17

2538

1.27 1.49 1.68 1.76 2.90 4.11 4.86 7.12

2233 2042 1821 1867 1517 1257 1326 1076

7% GAg47 μm 11% GAg47 μm 17% GAg47 μm 23% GAg47 μm 34% GAg47 μm 37% GAg47 μm 49% GAg47 μm

0.29 0.33 0.41 0.52 0.70 0.80 1.28

1.30 1.58 1.93 2.33 3.04 3.73 7.88

2209 1917 1778 1700 1542 1406 935

EVA (28/40) + Baryum Titanate = BT 5% BT9 μm 10% BT9 μm 16% BT9 μm 28% BT9 μm 38% BT9 μm 48% BT9 μm

0.27

1.17

2538

0.27 0.31 0.36 0.54 0.69 0.89

1.30 1.50 1.95 2.63 3.18 3.90

1781 1473 1061 886 766 681

Material

Specific heat capacity J.kg−1 .K−1

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Table 10.2 (Continued) Material

Thermal conductivity W.m−1 .K−1

Thermal diffusivity (m2 /s)· 10−7

Specific heat capacity J.kg−1 .K−1

5% BT105 μm 12% BT105 μm 20% BT105 μm 27% BT105 μm 36% BT105 μm 44% BT105 μm

0.30 0.36 0.48 0.58 0.74 0.90

1.38 1.79 2.20 2.66 3.16 4.03

1935 1278 1103 946 852 717

Polycarbonate 10% glass 30% glass

0.20 0.22 0.32

—

1255

Epoxy resin 30 % mica 50 % mica

0.16 0.24 0.39

— — —

1045

Polypropylene 40% talc 40% CaCO3 40% glass

0.12 0.32 0.29 0.37

— — — —

2350

—

According to Sundstrom [2], theoretical models are classified into two classes: complete models and simplified models. In the first case, the effective conductivity is obtained using an analytical solution of the heat equation without any assumptions about the flow of heat or material structure. In the case of simplified models, it is generally considered that the heat flow is unidirectional and that the isotherms are planes perpendicular to the heat flow. Mottram [57] proposed an alternative classification of theoretical models: models of first, second, third and fourth order. This last classification will be used in this chapter for the presentation of the most commonly-used prediction models. Models of First Order The simplest alternative for modeling thermal conductivity of a two-phase system is to represent the material using two components arranged in either parallel or series with respect to the heat flow (see Figure 10.11). This gives the upper (ksup ) and lower (klow ) bounds of effective thermal conductivity: parallel model series model

ksup = km (1 − ϕ) + k f ϕ (1 − ϕ) 1 ϕ = + klow km kf

(10.4) (10.5)

where kp and kf are the thermal conductivities of the polymeric matrix and the fillers respectively and ϕ is the filler volume fraction. This simple modeling provides upper and lower values for the thermal conductivity of a two-phase material. So, the exact value of the thermal conductivity is always located between these bounds. In most cases, estimations provided by a first order model are far from the value obtained from measurement. However, in some particular situation these models can provide quite good estimations of the thermal conductivity. For instance, in the case of the use of fibers arranged perpendicular to the heat flow with a ratio between

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405

Thermal Flux

k1 k2 klow

ksup

ki kN (a) Series conduction

Figure 10.11

k1 k2 k3 ki kN (b) Parallel conduction

Schematic representation of parallel conduction and series conduction.

fibers and matrix thermal conductivity values lower than 10, the lower bound of this model is useful for the prediction of the effective thermal conductivity. In polymer composites, if the alignment of the particles fillers or fibers is parallel to the direction of the thermal flux, the thermal conductivity will be the highest one (parallel conduction) and if the alignment is perpendicular to the direction of the thermal flux, thermal conductivity will be the lowest (series conduction). Second Order Models These models usually take into account the same parameters as the first order models. Some models of second order consider also additional parameters such as the shape of the dispersed phase (sphere, cylinder, fibers . . .). Maxwell’s approach, initially associated with a problem of electrical conduction in a heterogeneous material consisting of spheres dispersed in a matrix, has led to the formulation of effective conductivity of a composite as:

k = km

  2km + k f − 2ϕ km − k f   2km + k f + ϕ km − k f

(10.6)

The Maxwell equation cannot be applied at high concentration of particles as the interaction between the particles is not considered in the model. For composites including higher filler concentration, numerous relationships have been proposed in the literature [67]. The most frequently used are described in the following. The Bruggeman model was originally established for the estimation of the dielectric permittivity of heterogeneous materials [13]. This model has subsequently been widely used for predicting the thermal conductivity of composite materials: λf − λ 1−ϕ = λ f − λm



λm λ

 d1 (10.7)

The value of parameter d depends on the shape of the fillers used: d = 3 for spherical fillers and d = 2 for cylindrical fillers.

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The Hatta and Taya model was originally based on the analogy of thermal conductivity to mechanical modulus [13, 57, 68]. The general equation corresponding to this model is:       k f − km (2S33 + S11 ) + 3km km − k f (10.8) k = km 1 − ϕ  2   3 k f − km (1 − ϕ) S11 S33 + km k f − km R + 3km2 where R = 3 (S11 + S33 ) − ϕ (2S11 + S33 )

(10.9)

The parameters Sii obey S11 + S22 + S33 = 1, and depend on the shape of the fillers used:

r r r r

discs: S11 = S22 = 0, S33 = 1 spheres: S11 = S22 = S33 = 1/3 fibers, rods or long cylinders: S11 = S22 = 1/2, S33 = 0 randomly-oriented short fibers (length L and diameter D):    ⎧ 1/2 α 2 −1 ⎪ S α α = − 1 − cosh α ⎪ 11   3/2 ⎪ ⎪ 2 α2 − 1 ⎪ ⎨ S33 = 1 − 2S11 ⎪ ⎪ ⎪ ⎪ ⎪ ⎩α = L D

(10.10)

In the case of random dispersion of spheres in a continuous matrix, R = 2 − ϕ The solution of the Hatta and Taya model for spherical fillers is equivalent to the solution of lower bound of the Hashin and Shtrikman model [13, 57, 69]: ⎛ ⎞ ⎜ k = km ⎜ ⎝1 +

ϕ

km 1−ϕ + 3 k f − km

⎟ ⎟ ⎠

(10.11)

Figure 10.12 shows an example of a comparison between experimental data, Hatta and Taya and Bruggeman models predictions in the case of composites of ethylene-vinyl acetate (EVA) filled with BaTiO3 powder of mean diameter equal to 9 μm and 105 μm for various volume filler loading [65, 66]. In this particular case, a good agreement between experimental data and both models estimations is noted. The Meredith and Tobias model [70] allows obtaining effective thermal conductivity values intermediate between the values given by the Maxwell model and those obtained from the Bruggeman model (see Figure 10.13). The Meredith and Tobias model is expressed as follows:           (2 − ϕ) 2 + k f /km − 2 1 − k f /km ϕ 2 2 + k f /km − 2 1 − k f /km ϕ         k = km (2 − ϕ) 2 + k f /km + 1 − k f /km ϕ 2 2 + k f /km + 1 − k f /km ϕ

(10.12)

if k f /km → ∞, k = km

(1 + ϕ) (2 + ϕ) (1 − ϕ) (2 − ϕ)

(10.13)

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Thermophysical Properties of Multiphase Polymer Systems 1.2

Exp. Data EVA/BaTiO3 9µm

1.0

EVA/BaTiO3105µm Bruggeman model

0.8

-1

-1

k (W.m .K )

407

0.6 0.4 0.2

Hatta & Taya model

0

10

20 30 ϕVol (%)

40

50

Figure 10.12 Variation of measured effective thermal conductivity values and those of the Hatta & Taya and Bruggeman models for composites EVA/BaTiO3 with filler mean diameter of 9 μm and 105 μm based on [66].

Based on Tsao’s probabilistic model [72], Cheng and Vachon [60, 73] assumed a parabolic distribution of the discontinuous phase. The constants of the parabolic distribution were evaluated as a function of the discontinuous phase volume fraction. The effective thermal conductivity is given for the case kf > km : 1 1−G = + k km  km × log  km

1   1/2  H k f − km km + G k f − km 1/2 1/2    + G k f − km + G/2 H k f − km    1/2 1/2 + G k f − km − G/2 H k f − km 

(10.14)

Figure 10.13 Comparison of some prediction models with measured values of thermal conductivity of epoxy resin loaded with aluminum nitride composite. Reprinted from [71]. Copyright (2008) with permission from American Ceramic Society.

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where  G=

3ϕ 2



1/2 ,

H=

2 3ϕ

1/2 (10.15)

The effect of interfacial thermal contact resistance between filler and matrix may play a major role on thermophysical properties of composites. Hasselman and Johnson modified the original models of Rayleigh and Maxwell to obtain finally the expression of effective thermal conductivity of two-phase composites with a dilute dispersion with spherical, cylindrical and flat-plate geometry fillers and a thermal barrier at the interface between the components [57]. This work has opened the way to other studies to understand the mechanism of heat transfer on threecomponent composite materials containing core-shell particles randomly distributed in a continuous matrix. Particles of core-shell type consist of a homogeneous core surrounded by a homogeneous shell of a different material. The encapsulation of core with shell is useful in many situations: to improve the thermal and/or mechanical properties, to modify the surface properties of particles, and also to improve the dispersion of particles in the matrix. A specific model was presented by Benveniste and Miloh to predict the thermal conductivity of coated fillers (particles or fibers) [74]. In that model, composites are seen as an assembly of three distinct phases: index m-continuous phase, 2-center of dispersed phase, c-coating of the dispersed phase. This dilute suspension model for spherical (d = 3) or cylindrical (d = 2) inclusions is defined by: k = km

where

and

β¯ =

β2c

1 + (d − 1) (ϕ2 + ϕc ) β 1 − (ϕ2 + ϕc ) β¯ k −1 km

; k + (d − 1) km (k2 /kc ) − 1 = (k2 /kc ) + (d − 1)

  ϕ2 β2c 1 + (d − 1) k kc ϕ2 + ϕc   = · ϕ2 km km β2c 1− ϕ2 + ϕc

(10.16)

(10.17)

(10.18)

We note that ϕ = ϕ2 + ϕc is the volume fraction of the fillers (core-shell) and k is the effective thermal conductivity of coated fillers. Krupa et al. have studied the thermophysical properties of EVA copolymer filled with silver coated wollastonite fibers. In this study, the computations were performed using the experimental value of the EVA matrix thermal conductivity (km = 0.131 W.m−1 .K−1 ) and a value of the fillers thermal conductivity (k = 131 W.m−1 .K−1 ). It is seen in Figure 10.14 that the Hatta and Taya model shows a good correlation with the experimental data up to 40 vol.% of the filler. In contrast, the Benveniste and Miloh model underestimates the composites thermal conductivity except at the lowest volume fractions (<10%) [75]. Recently, Rajinder Pal developed new thermal conductivity prediction models derived using the differential effective medium approach for three-component composites of core-shell particles. According to this author, the relative thermal conductivity (k) of a particulate composite, defined as composite conductivity divided by matrix conductivity, is a function of five variables, namely: ratio of shell-to-core radii (δ); ratio of shellto-matrix thermal conductivities (k21 ); ratio of core-to-shell thermal conductivities (k32 ); volume fraction of core-shell particles (ϕ); and maximum packing volume fraction of particles (ϕ m ) [76]. The summary of the four models developed by Pal is presented in Table 10.3.

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409

2.0 Hatta & Taya Benveniste & Miloh Exp. Data

-1

-1

k (W.m .K )

1.5

1.0

0.5

0.0 0

10

20

30

40

Filler volume fraction (%)

Figure 10.14 Comparison of experimental thermal conductivity values of EVA/W-Ag composites to some prediction models.

Table 10.3 Summary of the new models developed by Rajinder Pal [76]. Model description   β −1 1/3 exp (ϕ) = k β −k

(1 − ϕ)−1 = k1/3

ϕ 1−



⎟ 1/3 ϕ ⎠=k



ϕmax ⎞−ϕmax

⎛ ⎜ ⎝

β −1 β −k

⎞

⎛ ⎜ exp ⎝



Comments

ϕ 1−

⎟ ϕ ⎠

ϕmax

β −1 β −k

 = k1/3



β −1 β −k



This model is valid for low to moderate values of ϕ (ϕ < 0.2). In its derivation, it is assumed that all the volume of the composite before new particles are added is available as free volume to the new particles. Thus, this model underestimates the crowding effect of particles [76]. This model is an improvement over model 1 as far as the effect of particle volume fraction (ϕ) is concerned. However, like model 1, this model does not consider the packing limit of particles. It ignores the fact that an upper limit exists for particle volume fraction. For random close packing of rigid uniform spheres, the maximum packing volume fraction of particles (ϕ max ) is 0.637. Thus, this model also underestimates the crowding effect of particles [76]. This model takes into account the packing limit of particles. It is derived using the approach of Mooney to take into account the crowding effect of particles. However, as pointed out by Krieger, Mooney’s approach tends to overestimate the crowding effect of particles [76]. This model covers a broad range of ϕ (0 ≤ ϕ ≤ ϕ max ) taking into consideration the packing limit of particles [76].

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3

φ = 0.45 φm = 0.637 δ = 4/3 λ32 = 10

4 2 1

10 Kr

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100 1000 1E4

1E5

λ21 (b) 100

Model 4 φm = 0.637

δ = 4/3

φ = 0.55

λ32 = 10

φ = 0.45

10

φ = 0.30

Kr

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1

10

100 1000 1E4

1E5

λ21

Figure 10.15 (a) Relative thermal conductivity (Kr ) of composite of core-shell particles predicted from the 4 models developed by Rajinder PAL. (b) Effect of filler volume fraction. Reprinted from [76]. Copyright (2008) with permission from Elsevier.

Parameter β in the equations shown in Table 10.3 is expressed as follows:     2 + δ 3 λ31 − 2 2 + δ 3 λ21    β=  1 + 2δ 3 λ31 − 1 + δ 3 λ32

(10.19)

where λ21 is the ratio of shell material thermal conductivity (k2 ) to matrix thermal conductivity (k1 ), λ31 is the ratio of core-to-matrix thermal conductivities (k3 /k1 ), and λ32 is the ratio of core-to-shell thermal conductivities (k3 /k2 ). Thus, λ31 = λ21 λ32 . Figure 10.15 shows the relative thermal conductivity predicted from the models presented in Table 10.3, under the conditions: δ = 4/3, λ32 = 10, ϕ = 0.45, ϕ max = 0.637. Third and Fourth Order Models These models take into account the disturbance between the phases of the composite and the geometry of the inclusions. They show an additional term based on a statistical distribution [32, 77, 78]. According to Torquato, strong enhancement of heat flow is obtained only when the maximum

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Table 10.4 Values of the three–point parameter ξ for the spherical inclusions. Filler volume fraction

Fully penetrable spheres

Random impenetrable spheres

0.10 0.20 0.30 0.40 0.50 0.60 0.65 0.70

0.056 0.114 0.171 0.230 0.290 0.351 −0.415

0.021 0.040 0.059 0.084 0.141 0.328 —

packing fraction is reached; heat transfer in composite materials is far from being comparable to electrical transport [57, 77, 79]. Torquato developed a model which is more general than that of Hashin and Shtrikman by introducing parameters that take into account the statistical perturbation around each particle, particularly a microstructural parameter derived using a set of statistical functions [80]: The expression of Torquato model is: kc 1 + φβ − (1 − φ) ζβ 2 = km 1 − φβ − (1 − φ) ζβ 2 k f − km β= k f + (d − 1) km

(10.20) (10.21)

where d is a constant depending on the dimensions of the system (d = 3 for spheres and 2 for cylinders) and ξ is the microstructural three-point parameter (0 < ξ < 1) and its corresponding values are given in Table 10.4 for different types of mono-dispersed spherical inclusions in the case of penetrable or impenetrable random distributions [57, 80, 81]. These different configurations of composite materials are shown in Figure 10.16. The values of the maximum packing fraction ϕ max are given in Tables 10.5 and 10.6. Figure 10.17 shows a comparison between the experimental values of thermal conductivity and some theoretical models (including the Torquato model) predictions for composites made of copper fillers dispersed

(a)

(b)

Figure 10.16 Dispersion of mono-dispersed spheres within a polymeric matrix: (a) full penetrable spheres; (b) random impenetrable spheres (based on [13, 57]).

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Table 10.5 Values of maximum packing fraction of common filler shapes (based on [57, 80]). Nature of filler

Shape of filler

Type of arrangement

ϕ max

Any Any Any Any Any Any Any Any Talc Mica wollastonite glass/carbon

Spheres Spheres Spheres Spheres/irregular Fibers, cylinders or rods Fibers, cylinders or rods Fibers, cylinders or rods Fibers, cylinders or rods Flakes Flakes Fibers Flakes

Face centred cubic or hexagonal Body centred cubic Simple cubic Random close Hexagonal close Simple cubic Random Random 3D Random 3D Random 3D Random 3D Random 3D, L/D = 10

0.7405 0.60 0.524 0.637 0.907 0.785 0.82 0.52 0.40–0.56 0.38–0.45 0.62 0.42

in a polypropylene matrix (PP/Cu) for two fillers size (samples loaded with small particles of 23 μm diameter (Cu(1) ) and samples loaded with large particles of 230 μm (Cu(2) )) [82]. It is found that the variation of thermal conductivity is in agreement with various theoretical models only for low concentrations. 10.4.2.2

Semi-empirical Prediction Models

Semi-empirical or semi-theoretical models use both theoretical laws and experimental results for predicting effective thermal conductivity. These models typically include adjustable parameters, which in most cases allow to fit the experimental results quite accurately. However, generally these adjustable parameters do not have physical meaning and must also be determined for each system studied. Mean Geometric Model One of the most widely-used semi-empirical models is the Ratcliffe Model (the name also meaning geometric model) [83]. Although this model does not use any physical law, it nevertheless sometimes gives a good estimation of the effective thermal conductivity of heterogeneous materials (Figure 10.18). The model defines the thermal conductivity of the blend or composite as the product Table 10.6 Values of Parameter A of the Lewis & Nielsen model for common filler types. Direction of heat flow

Dispersed phase

A

Any

1.50

Any Any Any Any Any Any

Spheres Random fibers, 3D: L/D = 2* L/D = 4 L/D = 6 L/D = 10 L/D = 15 L/D = 35

Parallel

Uniaxially oriented fibers

Perpendicular

Uniaxially oriented fibers

* L/D:

ratio aspect of fibers

1.58 2.08 2.8 4.93 8.38 29 L 2 D 0.5

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Thermophysical Properties of Multiphase Polymer Systems Torquato model Hatta & Taya model Bruggeman model Exp. Data (PP/Cu(1)) Exp. Data (PP/Cu(2))

2.5 2.0

-1

-1

k (W.m .K )

413

1.5 1.0 0.5 0

10

20 30 40 Filler volume fraction (%)

50

Figure 10.17 Comparison of experimental thermal conductivity values of PP/Cu composites to some prediction models based on [82].

of thermal conductivity of each phase weighted by its volume fraction: ϕ (1−ϕ f ) k = k f f .km

(10.22)

Agari Model Agari’s model [84] is also extensively used in most works devoted to the study of the thermal conductivity of polymer matrix composites. Agari proposed another approach for all types of inclusions in a matrix. The model can be used even for multiphase systems and is usable for both low and high loading. This model is based on the generalization of series and parallel conduction models (first-order model) in the k kf 1

1-Series model 2-Parallel model 3-Mean geometric model

0.8

0.6 2

3

0.4

1

0.2

Series 0

Figure 10.18

0

0.5

ϕ 1

Schematic representation of parallel, series and mean geometric conduction models.

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composite and correlates thermal conductivity with the ability of fillers to create particle conductive chains [84, 85]. The logarithmic equation of Agari is: log k = ϕC2 log k f + (1 − ϕ) log (C1 km )

(10.23)

where parameters C1 and C2 are obtained by fitting experimental data. According to Agari C1 represents the effect of particles on the polymer structure, i.e. C1 is related to the change of thermal conductivity of polymeric matrix, as a consequence of a change of its crystallinity; C1 = 1 if no change in crystallinity is observed or if an amorphous matrix is used; C2 represents the ability of filler particles to create continuous chains. Moreover, Agari and co-workers assumed that there is a relationship between C2 parameter obtained and the concentration at electrical percolation threshold in the case where metallic particles are used. Previous works showed that these two parameters are closely linked. This assumption is possible if the crystallinity of the polymer does not change in the presence of filler particles. In this case, parameter C1 = 1 and the Agari model becomes:     kf k = ϕC2 log (10.24) log km km Finally, Agari et al. showed that the C2 parameter remains a curve fitting parameter that can be used to determine the electrical percolation concentration threshold using [86–88]. 

1 C2 = log ϕc

 (10.25)

where ϕ c is determined as the volume concentration at the inflection point in the electrical resistivity versus filler concentration curve, i.e. concentration at electrical percolation threshold. Figure 10.19 presents a comparison between experimental data and predictions given from the Agari model in the case of a polystyrene (PS) matrix filled with Aluminium nitride (AlN) particles of diameter equal to 0.15 and 2 mm [89]. 0.4 log (Thermal Conductivity) (W/m.K)

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0.2 0 –0.2 –0.4 –0.6 –0.8 –1 0

10

20

30

40

Volume Fraction of AIN (%)

Figure 10.19 Comparison of thermal conductivity of the AIN-polystyren composites with Agari model. Reprinted from [89]. Copyright (2008) with permission from Elsevier.

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Lewis and Nielsen Model This semi-theoretical model is a modification of the Halpin–Tsai equation [83] originally developed for mechanical behavior. In 1974, Lewis extended the application of this model for the prediction of the effective thermal conductivity of two-phase systems by including the effect of the fillers shape and orientation and the effect of the fillers maximum packing fraction (ϕ max ) [90]. The proposed equation is: 1 + ABϕ 1 − Bψϕ (1 − ϕmax ) ϕ f ψ = 1+ 2 ϕmax k = km

where

(10.26) and

B=

k f /km − 1 k f /km + A

(10.27)

The constant A is related to Einstein coefficient λE, [90]: A = kE − 1

(10.28)

The constant A depends on the shape and orientation of the fillers. The most important values of of A and ϕ max for several fillers shape and orientation are given in Table 10.6. 10.4.2.3

Numerical Prediction Models

Numerical modeling constitutes another way to study complex systems. In some cases, the physical problems do not admit an analytically exact solution: numerical modeling is a useful tool in this situation. However, it is generally suitable to compare numerical results to experiments. A lot of methods that allow the solving of numerically physical problems have been developed in the past. At this time, an increasing amount of commercial software or freeware is available to perform numerical modeling. Numerous studies devoted to numerical prediction models of thermal conductivity of filler polymer composites can be found in the literature. The whole volume of a heterogeneous material sample cannot be considered for modeling due to limited performances and memory size of computers and to dramatic increase of computing duration. The general way for numerical modeling of thermophysical properties of heterogeneous materials like polymer composites, is to consider only a small part of the material as representative of its structure. Thus, macroscopic properties are obtained by considering a unit cell of the material containing a limited number of fillers. In this section we will focus the presentation only on a few works related to the development of numerical prediction models of thermal conductivity of polymer composites. Deissler’s [91] work, which can be considered as one of the oldest studies concerning numerical modeling, was extended by Wakao and Kato [92] for a cubic or orthorhombic array of uniform spheres in contact on a polymeric matrix. Shonnard and Whitaker [93] have investigated the influence of contacts on two-dimensional models. They have developed a global equation based on an integral method for modeling heat transfers in the medium. Using the finite elements method, Veyret et al. [94] studied the heat conductive transfer in a periodic distribution of fillers in composite materials. In their study, computation was carried out on twoand three-dimensional geometric spaces. The same method was used by Ramani and Vaidyanthan [95], who predicted the effective thermal conductivity of composites with circular and rectangular fillers considering two-dimensional (2D) models taking into account the effect of microstructural characteristics such as filler aspect ratio, interfacial thermal resistance, volume fraction, and filler dispersion. These authors showed that for 50 vol. % of spherical copper powder filler, the thermal conductivity of the composite is increased from

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0.32 W.m−1 .K−1 (for pure PA6 matrix) to 7.29 W.m−1 .K−1 . Similar studies were also reported by several authors [96–98]. Yin and Tu [99] developed a 2-D finite element (FEM) model applied to the estimation of effective thermal conductivity of PTFE filled with randomly distributed circular graphite particles. They identified a saturation value of the filler to matrix thermal conductivity ratio for some fillers volume fraction values [100]. Bakker calculated the thermal conductivity coefficient of porous materials by using FEM, and investigated the relationship between two-dimensional (2D) and three-dimensional (3D) values obtained [101]. Two-dimensional (2D) Numerical Modeling Effective thermal conductivity of high density polyethylene filled with aluminium particles was investigated by D. Kumlutas et al. using two-dimensional numerical

4.5 Numerical Experimental

4

Maxwell y = 0.5443e0.0601x R2 = 0.9894

Russel

3.5

Baschirow & Selenew Thermal Conductivity (W/mK)

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Lewis & Nielsen (A = 3, ϕ = 0.637) m

Cheng & Vachon 2.5

Expon. (Experimental)

2

1.5

1

0.5

0 0

10

20

30

40

Volumetric Ratio Al%

Figure 10.20 Comparison between analytical, semi empirical, 2D numerical model and experimental data of the HDPE filled Al particles. Reprinted from [98]. Copyright (2003) with permission from Elsevier.

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(a)

417

(b)

Figure 10.21 3D geometry used for 44% particle concentration (a) sphere in cube and (b) cube in cube. Reprinted from [102]. Copyright (2006) with permission from SAGE Publications.

models [98]. The temperature field in the composite material was obtained by solving numerically Laplace’s equation using a finite difference formulation by imposing boundary conditions. Figure 10.20 shows a comparison between experimental data of HDPE/Al composites thermal conductivity, 2D numerical model computation results and some analytical models. One can see that the analytical and semi-empirical models provide quite a good estimation of the thermal conductivity of this composite only for volume fractions of Al lower than 10%. However, it is seen that the results obtained using the 2D numerical model are closer to experimental values than analytical models [98]. Three-dimensional (3D) Numerical Modeling Using the finite element program ANSYS, Kumlutas and Tavman [102] estimated numerically using a 3D model the effective thermal conductivity of composite materials for various filler concentrations and different kf /km ratios. The unit cell considered for modeling was chosen as a cube and spherical or cubic shapes of fillers were considered (Figure 10.21). Also, using 3D model finite element program ANSYS, a numerical estimation of thermal conductivity of HDPE filled with tin particles was estimated [102]. The results were compared to analytical models and experimental data (see Figure 10.22). Cai et al. [103] studied the thermal conductivity of PTFE composites reinforced with one or two fillers (particles and fibers) using 3D finite element simulations with ANSYS software. These authors considered a unit cell of cubic shape and used a specific algorithm to generate positions and/or orientations of the fillers within the matrix for a given volume fraction [103]. The algorithm ensures that the fillers do not intersect by judging whether the new generated filler intersects the fillers that have already been created. Since there is no limitation in size and shape of charges in the build process, it is considerably easier to generate a set of 3D randomly distributed charges with different materials and shapes (Figure 10.23). Figure 10.24 shows a comparison between 3D finite element model results proposed by Cai et al., some theoretical model and experimental values in the case of PTFE composites filled with graphite particles [103]. At the same time, several configurations were considered and are compared with each other in the same figure: (i) 3D model with random distribution of particles size, (ii) 3D model with uniform size of particles, and (iii) 2D model. It is seen that the effective thermal conductivity of PTFE/Graphite composites predicted by the 3D model with randomly distribution size of particles is in better agreement with experimental data than all other models.

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Handbook of Multiphase Polymer Systems 1.4 Numerical 1.3

Experimental Cheng and Vachon

1.2

Effective thermal conductivity (W/mK)

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Lewis and Nielsen (A = 3, ϕm = 0.637) Agari and Uno

1.1

Lewis and Nielsen (A = 1.5, ϕm = 0.637) Maxwell

1

0.9

0.8

0.7

0.6

0.5

0

5

10

15

20

Sn%

Figure 10.22 Comparison of a 3D sphere in cube numerical results, models and experimental data of tin filled HDPE at 38 ◦ C. (C1 = 0.938 and C2 = 0.897 of Agari’s model). Reprinted from [102]. Copyright (2006) with permission from SAGE Publications.

y

(a)

z

x

(b)

Figure 10.23 The random 3D finite element model of representative volume element size distribution (A) spheres with various diameters, (b) fibers. Reprinted from [103]. Copyright (2005) with permission from SAGE Publications.

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1.0 Effective thermal conductivity (W/mK)

419

0.9 0.8

2-D FE Model Maxwell Russell

0.7

Cheng and Vachon Nielsen

0.6 0.5 0.4 0.3 0.2 0

5

10

20 15 Volume fraction (%)

25

30

Figure 10.24 Comparison of the finite element model, theoretical models and experimental values of PTFE composite filled with graphite particles. Reprinted from [103]. Copyright (2005) with permission from SAGE Publications.

10.5 Summary The thermophysical properties of polymers, and more particularly thermal conductivity, can be enhanced by orders of magnitude using high thermal conductivity of fillers such as metal or metal oxides, boron oxide, carbon graphite . . . A wide variety of polymer materials such as thermoplastics, thermosets and elastomers can be used as the matrix material. Thermally conductive polymer composites find wide applications in several engineering fields, for instance in electronic components packaging or the heat dissipation of circuit boards. In aeronautics and the space domain, many applications require at the same time light weight and good thermal conduction. Thermal conductivity of composite materials is governed by many parameters, such as the concentration of fillers in the matrix, the size and shape of inclusions, the interactions between particles and the matrix . . . The diversity of these parameters and the degree of influence of each parameter are the root of the difficulty of reaching accurate modeling of thermophysical properties. However, most prediction models are able to give rather good estimations of thermal conductivity for low and medium conductive filler volume concentration corresponding to the electrical percolation threshold. This means that heat transfer in composites including conductive fillers behaves differently below and above the electrical percolation threshold. Thus, even if a great amount of work was devoted in the past to modeling thermophysical properties of heterogeneous materials, the development of prediction models is still necessary. In the same way, development of new materials with specific thermophysical properties still requires experimental characterization studies to be performed. At the start of this chapter, we showed that a lot of characterization techniques are now available to obtain experimental values of thermophysical parameters of polymers, blends or composites. Some techniques also allow obtaining experimental values as a function of temperature and sometimes in solid state as well

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as in liquid state. Recent and further developments of characterization techniques should follow different objectives:

r r r r r

characterization of materials with orthotropic thermophysical properties, for instance, carbon fiberreinforced plastics whose applications are numerous in transport and civil engineering; multi-scale characterization of composites since effective models require parameter values that are still difficult to measure like, for instance, thermal properties of anisotropic particles, thermal contact resistance between matrix and fillers; characterization of high-performance thermal insulation materials which are highly porous materials; application of existing techniques to on-line characterization of materials during manufacturing stage and to in situ characterization of materials; characterization of composite materials during their use, for instance by performing in situ periodic thermal characterizations: this can be suitable for early detection of defect appearance like, for instance, delaminating of fiber-reinforced composites.

References 1. Bantrup, J., Immergut, E., Grulke, E. Polymer Handbook, 4th ed. New York, NY: John Wiley & Sons, Inc.; 1999. 2. Sundstrom, W., Lee, Y.D. Thermal conductivity of polymers filled with particulate solids, Journal of Applied Polymer Science, 16, 3159–3167, 1972. 3. Bujard, P., Munk, K., Kuehnlein, G. In: Tong TW (Ed.) Thermal Conductivity, Vol. 22. Lancaster: Technomic Publishing Co. Inc., 1994. 4. Krupa, I., Choda,’K. I. Eur Polym J., 37(11): 2159–68, 2001. 5. Mamunya, YeP., Davydenko, V.V., Pissis, P., Lebedev, E.V. Eur Polym J, 38(9), 1887–97, 2002. 6. Budiansky, B. J. Comp. Mat. 14, 120, 1980. 7. Nomura, S., Chou, T.W. J. Comp. Mat. 14, 120, 1980. 8. Ma Chen-Chi, M., Tseng Han-Thing, H.D., Wu, J. Appl. Polym. Sci. 68, 1119, 1998. 9. Zalba, B, Marı’n, J.M., Cabeza, L.F., Mehling, H. Appl Therm Eng, 23(3), 251–83, 2003. 10. Boudenne, A., Ibos, L., Fois, M., Gehin, E., Majeste, J.C. Electrical and thermal behavior of polypropylene filled with copper particles, Composites: Part A, Appl Sci Manufact., 36, 1545–1554, 2005. 11. Kusy, R.P, Corneliussen, R.D. The thermal conductivity of nickel and copper dispersed in poly(vinyl chloride). Polym Eng Sci., 15(2), 107–112, 1975. 12. Young Seok Dong, Jae Rayoun Youn. Influence of dispersion states of carbon nanotubes on physical properties of epoxy nanocomposites. Carbon, 43, 1378–1385, 2005. 13. Godovsky, Y.K., Privalko, V.P. Thermal and Electrical Conductivity of Polymer Materials. Springer-Verlag: Berlin, 1995; pp. 1–30. 14. S¨oderstr¨on, T., Stoica, P. System Identification. Prentice Hall: Upper Saddle River, NJ, 1989. 15. Akahane, T., Kondoh, M., Hashimoto, K., Nagakawa, M. Japan. J. Appl. Phys., 26, 1000, 1987. 16. Preethy Menon, C., Philip, T. Meas. Sci. Technol., 11, 1744, 2000. 17. Sun, J., Longtin, J.P., Irvine, T.F. Int. J. Heat. Mass Transfer, 44, 645, 2001. 18. Salmon, D. Thermal conductivity of insulations using guarded hot plates, including recent developments and sources of reference, Materials, Meas. Sci. Technol., 12, 89–98, 2001. 19. Cahill, D.G., Pohl, R.O. Thermal conductivity of amorphous solids above the plateau, Phys. Rev. B 35, 4067, 1987. 20. Ohmura, T., Tsuboi, M., Tomimura, T. International Journal of Thermophysics, Vol. 23, No. 3, 2002. 21. Dowding, K.J., Beck, J.V., Blackwell, B.F. International Journal of Heat and Mass Transfer, 39: 15, pp. 3157–64, 1996. 22. Sawaf, B., Ozisik, M.N., Jarny, Y. International Journal of Heat and Mass Transfer, 38: 16, 3005–10, 1995. 23. Assaf, B., Menge, H., Sobotka, V., Trochu, F. Journal of Reinforced Plastics and Composites, 24, 1837, 2005.

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24. Degiovanni, A., Batsale, J.C., Maillet, D. Revue Generale de Thermique, 35 :141–147, 1996. 25. Enguehard, F., Bosher, D., D´eom, A., Balageas, D. Materials Science and Engineering B, 5, 127–134, 1990. 26. Parker, W.J., Jenkins, R.J., Butler, C.P., Abbott, G.L. Flash method of determining thermal diffusivity, heat capacity and thermal conductivity, J. App. Phys., 32: 9, 1679–1684, 961. 27. Baba, T., Ono, A. Improvement of the laser flash method to reduce uncertainty in thermal diffusivity measurements, Meas. Sci. Technol., 12, 2046–2057, 2001. 28. Degiovanni, A. Identification de la diffusivit´e thermique par l’utilisation des moments temporels partiels. High Temp., High Pressure, 17, p. 683, 1985. 29. Weig, G., Zhang, X., Yu, F., Chen, Kui, Thermal diffusivity measurements on insulation materials with the laser flash method, International Journal of Thermophysics, 27: 1, January 2006. 30. Cheng, Stephen Z.D. Handbook of Thermal Analysis and Calorimetry, Volume 3, Applications to Polymers and Plastics, Elsevier Science B.V, 2002. 31. Wunderlich, B. Thermal Analysis, Academic Press, Boston, 1990. 32. Thompson, E.V. Thermal properties, Encyclopedia of Polymer Science and Engineering, Volume 16, Second Edition, John Wiley & Sons Inc., New York, NY, 1989, pp. 771–747. 33. Balageas, Daniel, Ory, Dominique. Generalized equation of touchau-type effusimeter; improvement of the measuring method for thermal effusivity (thermal inertia)Verallgemeinerte gleichung des touchau-effusivit¨atsmessers. International Journal of Heat and Mass Transfer, 23: 3, 339–347, March 1980. 34. Zammit, U., Marinelli, M., Pizzoferato, R., Scudieri, F., Martellucci, S. J. Phys. E: Sci. Instrum., 21, 935, 1988. 35. Garnier, B., Delaunay, D., Beck, J.V. Int. J. Thermophys., 13, 1097, 1992. 36. Watanabe, H. Metrologia, 33, 101, 1996. 37. Thoen, J., Glorieux, C. Thermochimica Acta., 304, 137, 1997. 38. Marinelli, M., Zammit, U., Mercurit, F., Pizzoferrato, R. J. Appl. Phys., 72(3), 1096, 1992. 39. Gustafsson, S.E. Rev. Sci. Instrum., 62, 797, 1991. 40. Gustafsson, S.E., Karawacki, E., Khan, M.N. J. Phys. D: Appl. Phys., 12, 1411, 1979. 41. Krapez, J.C. Mesure de l’effusivit´e thermique – M´ethode par contact, Techniques de l’Ing´enieur, Trait´e Mesures et Contrˆole R, 2958, 2007. 42. Pettersson, S. Rev. Sci. Instrum., 61, 1308, 1990. 43. Karawacki, E., Suleiman, B.M. Meas. Sci. Technol., 2, 744, 1991. 44. Nunes Dos Santos, Wilson, Gregorio Rinaldo JR, Hot-wire parallel technique: A new method for simultaneous determination of thermal properties of polymers, Journal of Applied Polymer Science, 85, 1779–1786, 2002. 45. Boudenne, A., Ibos, L., Gehin, E., Candau, Y. Journal of Physics D: Applied Physics, 37, 132–139, 2004. 46. Boudenne, A., Ibos, L., Candau, Y. Measurement Science and Technology, 17, 1870–1876, 2006. 47. Boudenne, A., Ibos, L., Datcu, S., Candau, Y. Thermal Conductivity 29: Thermal Expansion17, Conference Information: Joint 29th International Thermal Conductivity Conference/17th International Thermal Expansion Symposium, June 24–27 2007, SE Res Inst, Birmingham, AL 437–446, 2008. 48. Agoudjil, B., Datcu, S., Boudenne, A., Ibos, L., Candau, Y. Review of Scientific Instruments, 77, 35106–35109, 2006. 49. Finlayson, D.M., Mason, P. J. Phys. C: Solid State Phys., 18, 1777, 1985. 50. Pattnaik, S., Thompson, E.V. Polym Prepr., 22(1), 299, 1981. 51. Choy, L.C., Leung, W.P., Ng, Y.K. J. Polym. Sci. Polym. Phys. Ed., 25. 1779, 1987. 52. Mark, James E. Physical Properties of Polymers Handbook, 2nd Ed., New-York: Springer, 2007. 53. Yang, Y. Thermal conductivity. In: J.E. Mark (Ed.), Physical Properties of Polymers Handbook, American Institute of Physics, Woodbury, NY, pp. 155–163, 1996. 54. Osswald, Tim A., Hernandez-Ortiz, Juan P. Polymer Processing – Modeling and Simulation, Hanser Gardner Publications, 2006. 55. Zhang, X., Fujii, M. Measurements of the thermal conductivity and thermal diffusivity of polymers, Polym Eng Sci., 43(11), pp. 1755–1764, 2003. 56. Wen, J. Heat capacity of polymers. In: J.E. Mark (Ed.), Physical Properties of Polymers Handbook, American Institute of Physics, Woodbury, NY, pp. 145–154, 1996.

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57. Mottram, J.T., Taylor, R. Thermal transport properties. International Encyclopedia of Composite, Vol. 5. New York: VCH, pp. 476–96, 1991. 58. Gorring, R.L., Churchill, S.W. Thermal conductivity of heterogeneous materials, Chem. Eng. Progr., 57(7), 53–59, 1961. 59. Godbee, H.W., Ziegler, Z.T. Thermal conductivities of MgO, Al2 O3, and ZrO2 powders to 850 ◦ C, J. Appl. Phys., 37, 56–65, 1966. 60. Cheng, S.C., Vachon, R.I. A technique for predicting the thermal conductivity of suspensions, emulsions and porous materials, Int. J. Heat Mass Transfer, 13, 537–546, 1970. 61. Tlili, R., Boudenne, A., Cecen, V., Ibos, L., Krupa, I., Candau, Y. Thermophysical and electrical properties of nanocomposites based on ethylene-vinylacetate copolymer (EVA) filled with expanded and unexpanded graphite, International Journal of Thermophysics, 31: 4–5, 936–948 (13), 2010. 62. Boudenne, A., Ibos, L., Fois, M., Gehin, E., Majeste, J.C. Thermophysical properties of (Polypropylene/Aluminum) composite materials, Journal of Polymer Science: Part B: Polymer Physics, 42, 722–732, 2004. 63. Krupa, I., Boudenne, A., Ibos, L. Thermophysical properties of polyethylene filled with metal coated polyamide particles, European Polymer Journal, 43, 2443–2452, 2007. 64. Agoudjil, B., Ibos, L., Candau, Y., Majest´e, J.C. Mamunya, Ye.P. Correlation between transport properties of Ethylene Vinyl Acetate/Glass, Silver-Coated Glass Spheres composites. Composites Part A, 39, 342–35, 2008. 65. Agoudjil, B., Ibos, L., Candau, Y., Majest´e, J.C. Dielectric and thermophysical properties of Ethylene Vinyl Acetate/BaTiO3 composites. IEEE International Conference on Solid Dielectrics, 1, 2, 385–388, 2007. 66. Agoudjil, B., Ibos, L., Candau, Y., Majest´e, J.C. Dielectric and thermophysical properties of Ethylene Vinyl Acetate/BaTiO3 composites, Journal of Physics D: Applied Physics, 41, 055407, 2008. 67. Pal, Rajinder. New models for thermal conductivity of particulate composites, Journal of Reinforced Plastics and Composites, 26, 643, 2007. 68. Hatta, H., Taya, M. Effective thermal conductivity of a misoriented short fiber composite. J Appl Phys., 58(7): 2478–86, 1985. 69. Hashin, Z., Shtrikman, S. A variational approach to the theory of the effective magnetic permeability of multiphase materials, J Appl Phys., 33(10), 3125–31, 1962. 70. Meredith, R.W., Tobias, C.W. Conduction in heterogeneous systems. In C.W. Tobias (Ed.) Advances in Electrochemistry and Electrochemical Engineering, Vol. 2. John Wiley & Sons Inc., New York, 1962, p. 15. 71. Lee, S., Lee, S.-M., Shanefield, D.J., Cannon, W.R. Enhanced thermal conductivity of polymer matrix composite via high, Journal of the American Ceramic Society, 91: 4, 1169–1174, 2008. 72. Tsao, G.T. Thermal conductivity of two-phase materials, Industrial and Engineering Chemistry, 53, 395–397, 1961. 73. Cheng, S.C., Vachon, R.I. Int. J. Heat. Mass. Transfer, 12, 249, 1969. 74. Benveniste, Y., Lloh, T. On the effective thermal conductivity of coated short-fiber composites, Journal of Applied Physics, 69(3), 1337–1344, 1991. 75. Krupa, I., Cecen, V., Tlili, R., Boudenne, A., Ibos, L. Thermophysical properties of ethylene-vinyl acetate copolymer (EVA) filled with wollastonite fibers coated by silver, European Polymer Journal, 44(11), 3817–3826, 2008. 76. Pal, Rajinder. Thermal conductivity of three-component composites of core-shell particles; Materials Science and Engineering: A, 498(1–2), 135–141, 2008. 77. Torquato, S., Gstell, J., Chem., 78, 3262, 1983. 78. Bigg, D.M. Thermal and electrical conductivity of polymer materials. Adv Polym Sci., 119: 1–30, 1995. 79. Torquato, S. J. Mec. Phys. Solids, 45, 1421, 1997. 80. Torquato, S. J Appl Phys., 58, 3790–3797, 1985. 81. Boudenne, A., Khaldi, S. Temperature and liquid crystal concentration effect on thermal conductivity of poly(styrene) dispersed 5CB liquid crystal, Journal of Applied Polymer Science, 89, 481–486, 2003. 82. Boudenne, A., Ibos, L., Fois, M., G´ehin, E., Majest´e, J.-C., Candau, Y. Comparison of the thermal properties of polypropylene/copper composites to different theoretical or semi-empirical models, Physical and Chemical News, 29, 56–61, 2006. 83. Progelhof, R.C., Throne, J.L., Ruetsch, R.R. Methods for predicting the thermal conductivity of composite systems: A review, Polym. Eng. Sci., 16, 615–619, 1976.

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84. Agari, Y., Uno, T. J. Appl. Polym. Sci., 32, 5705, 1986. 85. Agari, Y., Tanaka, M., Nagai, S. J. Appl. Polym. Sci., 34, 1429, 1987. 86. Agari, Y., Ueda, A., Tanaka, M., Nagai, S. Thermal conductivity of a polymer filled with particles in the wide range from low to super-high volume content. J Appl Polym Sci., 40, 929–941, 1990. 87. Agari, Y,, Ueda, A., Nagai, S. Thermal conductivity of a polyethylene filled with disoriented short-cut carbon fibers. J Appl Polym Sci., 43, 1117–1124, 1991. 88. Agari, Y., Ueda, A., Nagai, S. Thermal conductivity of a polymer composite. J Appl Polym Sci., 49, 1625–1634, 1993. 89. Yu, S., Hing, P., Hu, X. Thermal conductivity of polystyrene–aluminum nitride composites, Composites Part A: Applied Science and Manufacturing, 33, 289–292, 2002. 90. Nielsen, L.E. The thermal and electrical conductivity of two phase systems, Ind. Engng Chem. Fund., 13, 17–18, 1974. 91. Deissler, R.G., Boegli, J.S. Trans ASME., 80, 1417, 1958. 92. Wakao, N., Kato, K. Chemical Engineering of Japan, 2, 24, 1969. 93. Shonnard, D.R., Whitaker, S., Int. J. Heat Mass Transfer, 32, 503, 1989. 94. Veyret, D., Cioulachtjian, S., Tadrist, L., Pantaloni, J. Journal of Heat Transfer, 115, 866, 1993. 95. Ramani, K., Vaidyanathan, A., Journal of Composite Materials, 29, 1725, 1995. 96. Islam, M.R., Pramila, A. Thermal conductivity of fiber reinforced composites by FEM, Journal of Composite Materials, 33, 1699–1715, 1999. 97. Jiang, M., Ostoja-Starzewski, M., Jasiuk, I. Apparent thermal conductivity of periodic two-dimensional composites, Journal of Computational Materials Science, 25, 329–338, 2002. 98. Kumlutas, D., Tavman, I.H., Coban, M.T. Thermal conductivity of particle filled polyethylene composite materials, Journal of Composites Science and Technology, 63, 113–117, 2003. 99. Yin, Y., Tu, S.T. Thermal conductivity of PTFE composites with random distributed graphite particles, Journal of Reinforced Plastic and Composites, 21, 1619–1627, 2002. 100. Tu, S.T., Yin, Y., Cai, W.Z., Ling, X. Saturation of physical properties of composites with randomly distributed second phase. In: MESO MECHANICS 2003, Tokyo, Japan, 26–28, 2003. 101. Bakker, K. Int. J. Heat Mass Transfer, 40: 15, 3503, 1997. 102. Kumlutas, D., Tavman, I.H. A numerical and experimental study on thermal conductivity of particle filled polymer composites, Journal of Thermoplastic Composite Materials, 19(4), 441–455, 2006. 103. Cai, W.-Z., Tu, S.-T., Tao, G.-L. Thermal conductivity of PTFE composites with three-dimensional randomly distributed fillers, Journal of Thermoplastic Composite Materials, 18(3), 241–253, 2005.

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11 Electrically Conductive Polymeric Composites and Nanocomposites Igor Krupa Polymer Institute, Slovak Academy of Sciences, D´ubravsk´a, Bratislava, Slovakia

Jan Prokeˇs Department of Macromolecular Physics, Faculty of Mathematics and Physics, Charles University in Prague, Prague, Czech Republic

Ivo Kˇrivka Department of Macromolecular Physics, Faculty of Mathematics and Physics, Charles University in Prague, Prague, Czech Republic

ˇ Zdeno Spitalsk´ y Polymer Institute, Slovak Academy of Sciences, D´ubravsk´a, Bratislava, Slovakia

11.1 Introduction Electrical properties of materials, in general, are characterized by their response when an electric field is applied to them. There are significant differences between the response of polymers and that of metals. In contrast to metals, in which the electrical field causes an electrical transport phenomena associated with a transport of electric charges, polymers may respond in a more varied manner and a whole set of electrical effects may be observed. Basically, the response of any material to an electric field can be separated into two main parts: (i) dielectric response, and (ii) electrical conduction. In frequency domain, the dielectric properties are characterized by the complex permittivity, the real part of which is often expressed in terms of the dielectric constant representing polarization effects. The imaginary part represents the dielectric losses related to relaxation phenomena.

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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The ratio of the imaginary to the real part of the complex permittivity is known as the tangent of dielectric loss angle. It should be noticed that, in spite of the name, the dielectric constant cannot be frequency independent over the whole frequency range. This is a consequence of the Kramers-Kronig relation between the real and imaginary parts of complex permittivity. Moreover, it must be kept in mind that the measured response always contains both the parts mentioned above because no device is able to separate the dielectric losses from the conductivity term. It is not a serious problem in the case of inorganic dielectrics whose conductivity can usually be omitted, and in the case of metals having negligible dielectric response. In contrast, the response of conducting polymers or polymeric composites contains simultaneously the conduction due to free charge carriers and the polarization of dipoles. Moreover, we cannot neglect the polarization due to charge carriers confined in isolated conducting clusters [1]. For practical purposes, all the materials are divided into three groups according to their electrical conductivity: (i) insulators having the electrical conductivity of the order of 10−8 S/m and lower; (ii) metals having the electrical conductivity of the order of 106 S/m and higher; (iii) the range between insulators and metals covers a large variety of materials – particularly semiconductors and semimetals, conducting polymers (e.g. doped polypyrrol, polyaniline, polyacetylene, polythiophene), composites and other heterogeneous materials [2]. Common non-conductive polymers are inherently insulating materials having an electrical conductivity of the order of 10−13 – 10−15 S.cm−1 [2]. For this reason, they are excellent electrical insulators. However, in many cases, it is desirable to have materials with some level of electrical conductivity, which also possess the unique properties of polymeric materials, such as low weight, easy and low cost processing, as well as various mechanical, thermal and optical properties, etc. [3, 4]. In order to obtain materials with the desired electrical, thermal and mechanical properties, polymers are frequently blended with different kinds of fillers. In dependence on the conductive filler content, these materials – called electrically conductive polymeric composites or nanocomposites (if at least one dimension of the filler has nano size) – may achieve the conductivity comparable to that of doped semiconductors. In this chapter we will discuss various aspects of these composites and nanocomposites.

11.2 Theory High performance conductive materials are formed by the incorporation of electrically conductive fillers into the polymeric matrix. Electrically conductive composites consisting of an insulating polymeric matrix and electrically conductive filler demonstrate typical sigmoidal behavior schematically shown in Figure 11.1. The percolation effect is observed in the dependence of conductivity versus filler content and manifests itself as a dramatic increase in conductivity by several orders of magnitude in a rather narrow concentration range of the filler around the so-called percolation threshold. In general, the percolation effect is a well-known phenomenon observed in filler-matrix systems as the abrupt extreme change of certain physical properties within a rather narrow concentration range of heterogeneity [3, 4]. The effect is explained as the formation of conductive paths through the matrix in such a way that the conductive particles are in close contact at a filler concentration corresponding to the percolation threshold. It was found that the formation of an internal network of filler within the matrix had a significant influence on other properties of the composites, like thermal conductivity [5], viscoelasticity [6] or some mechanical properties [7, 8].

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Electrically Conductive Polymeric Composites and Nanocomposites

I. insulating region

427

II. conductive region

percolation concentration

volume portion of the filler

Figure 11.1

11.2.1

Typical behavior of the dependence electrical conductivity versus filler content.

Percolation Models

Percolation theory was introduced in 1957 [9, 10] by Hammersley as a mathematical framework for the study of random physical processes such as flow through a disordered porous medium. Percolation theory deals with a generalized flow in random media. Electrical analogs include the flow of electricity through random resistor networks. If the medium is a set of regular lattice points, then there are two types of percolation. A site percolation considers the lattice vertices as the relevant entities; a bond percolation considers the lattice edges as the relevant entities. The terms ‘site percolation’ and ‘bond percolation’ [11] are schematically demonstrated in Figure 11.2 (a,b). It has been shown [12] that each bond percolation task can be translated to a site percolation one, but the opposite statement does not hold. The fundamental difference between site and bond percolation consists in the correlation between bonds connecting a particular site. Percolation threshold is a mathematical term relating to percolation theory, which is the formation of longrange connectivity in random systems. In engineering, the percolation is the slow flow of fluids through porous media or the current flow though a heterogeneous conductor, for instance, but in mathematics and physics it generally refers to simplified lattice models of random systems, and the nature of the connectivity in them.

bond percolation

Figure 11.2

site percolation

(a) Bond percolation, (b) site percolation.

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An important task is to find the so-called percolation threshold. The percolation threshold is the critical value of the occupation probability p, such that nonzero connectivity (percolation) first occurs. To describe electrical behavior of composites, a large number of different models were developed, taking into account various parameters. In the following section we will briefly discuss some of the most referred models. 11.2.1.1

Statistical Percolation Models

Statistical percolation models are the most studied and applied models for a characterization of electrical percolation phenomenon. The basis for that type of model was given in pioneering works by Kirkpatrick and Zallen [13, 14]. An estimation of percolation concentration is based on the random filling of finite regular array by points or bonds (bonds represent a connection between the points). Typical geometries of such arrays are the simple cubic lattice, the face-centered lattice, body-centred cubic lattice, etc. The calculation is focused on the determination of the part of existing points or bonds, incorporated in a big cluster within a studied array. The output of the calculation is the percolation (threshold) concentration of points ( pcsite ) or bonds ( pcbond ). In this case at least one infinity cluster penetrates through the array (it spans from one side of the array to the other). Table 11.1 summarizes threshold values pcsite and pcbond calculated for lattices of various type and dimensionality. In order to obtain blending equations (blending curves) for the description of the dependence of conductivity versus volume filler content, it is necessary to correlate the pc values with the conductivity of the mixture. Usually it is supposed that (at least in the vicinity of the critical concentration) the percolation equation [13, 14] can be used: σc = σ0 (φ f − φc )t

(11.1)

Table 11.1 Parameters related to the percolation theory of electrical conductivity of binary mixtures according to Zallen [14]. Dimensionality

Lattice or structure

pcbond

pcsite

Filling factor ν

νpcsite = φc

1 2 2 2 2

chain triangular square kagome honeycomb

1 0.3473 0.5000 0.45 0.6527

1 0.5000 0.593 0.6527 0.698

1 0.9069 0.7854 0.6802 0.6046

3 3 3 3 3

fcc bcc sc diamond rcp

0.119 0.179 0.247 0.388

0.198 0.245 0.311 0.428 0.27

0.7405 0.6802 0.5236 0.3401 0.637

4 4 5 5 6

sc fcc sc fcc sc

0.160

0.197 0.098 0.141 0.054 0.107

0.3804 0.6169 0.1645 0.4653 0.0807

1 0.45 0.47 0.44 0.42 0.45±0.03 0.147 0.167 0.163 0.146 0.16 0.16±0.02 0.061 0.060 0.023 0.025 0.009

0.118 0.094

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where σ c is the conductivity of the mixture (composite), σ 0 is the composite parameter proportional to the conductivity of the filler σ f (σ f = σ0 (1 − φc )t ) and takes into account any contact resistance between particles, φ f is the filler volume fraction, φ c is the volume percolation threshold (critical volume fraction; CVF) and t is the percolation exponent characterizing the power of the conductivity increase above percolation concentration. Kirkpatrick [13] computed the following values for the percolation coefficient t: t = 1.6 ± 0.1 for the bond percolation model and t = 1.5 ± 0.1 for site (or point) percolation model. Later calculations shifted the critical coefficient for site percolation to the value 1.9 ± 0.1 [15]. If we want to use Eq. (11.1) for a calculation of CVF values which are normally determined from experimental studies, we have to recalculate pc values using some geometrical considerations. Very rough simplification is adopted in [14], where the system is supposed to be composed of a disordered mixture of conducting (metallic) spheres of radius RM and insulating (or missing) spheres of radius RI , where (i) the conductive and isolating particles are of the same size (RM ∼ RI ); and (ii) the whole space between particles belongs to the isolating phase. The CVF value φ c can be then computed according to the relation: φc = ν

pcsite

(11.2)

where ν is the filling factor of the representing unit cell of a specific point arrangement. It can be seen from Table 11.1 that the CVF values φ c obtained using Eq. (11.2) are almost independent on the lattice type. Although based on very special assumptions, this result sometimes serves as a hint that CVF and critical exponent t are dimensional invariants, i.e. they only depend on the dimensionality of the system. Nevertheless, even for regular spherical particles, it can be easily demonstrated that CVF can achieve any value between 0 and 64 vol. %. In the case RM  RI , the conducting particles are embedded in an insulating matrix formed by fine particles filling the spaces in between. Such a system exhibits percolation only in the trivial limit of random close packing equal approximately to 64% of the total volume [16]. In mixtures fulfilling the condition RM  RI , the metal particles fill nonspherical voids between the insulating particles (segregated structure). Extensive analysis of this type of system [17] shows that CVF of metal may be reduced from 35 vol. % to 6 vol. % by a change of the ratio RI /RM from 1 to 16. If the conducting component can be considered rather as a continuum, then the ratio RI /RM → ∞ and consequently CVF → 0 [18–21]. Experimental measurements of blending curves for various composite systems also provide values of the critical exponent t. In spite of theoretical assumptions, a large variety of values can be observed [22]. Another frequently-applied model is that of Janzen [23]. This model is expressed by Equation (11.3) and enables to estimate the percolation concentration (φ c ) on the base of filler properties such as the bulk density of the filler (ρ) and the specific pore volume of the filler particles (ε). Parameter ε is usually determined experimentally from the absorption of some liquids (dibutylpthalate-DBP, for example): c =

1 1 + 0.67zρε

(11.3)

For commercial carbon black the DBP-absorption is between 0.9–3.8 cm3 /g. The term 0.67z is often considered equal to 4. With these parameters the theoretical threshold is ca. 3 vol. % or more [24]. As has been demonstrated above, the simple classical statistical models based on Eq. (11.1) can be applied to a large set of experimental results. However, a serious disagreement between the predictions

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of the simple models and experimentally found percolation behavior is often observed. For this reason, more complex percolation models were developed taking into account various preconditions. For example, Sumita and co-workers [25, 26] and Wessling and co-workers [27, 28] proposed thermodynamic percolation model, preferably applied for polymers filled with carbon black and with intrinsically conductive polymers. These models emphasize the importance of the interfacial interactions at the boundary between polymer and filler for the network formation. These models interpret the percolation phenomenon as a phase separation process. A further type of percolation model is represented by geometrical models. This kind of model is particularly applied to explain the percolation phenomenon in different dry-premixed and subsequently sintered mixtures of conductive and insulating powders. The best-known models of this class are those of Slupkowski [29], Rajagopal and Satyam [30], Malliaris and Turner [17], and Bhattacharya and Chaklader [31]. The details concerning those models can be found in the original papers, or briefly described and discussed in the comprehensive review by Lux [3]. Geometrical models can be applied not only to sintered materials but also to a certain type of polymeric composite [32]. Finally, the last group of percolation models are structural-oriented models. It seems to be generally accepted that an appropriate reproducibility in percolation concentration determination can only be found when the conductive composites are prepared via the hot melt blending followed by compression molding route [3]. This preparation technique ensures an appropriate homogeneity in distribution of the filler. Taking into account this fact, Yoshida [33], McCullough [34], Ondracek and co-workers [35], and Nielsen [36] proposed conductivity models on the basis of parameters, which can be determined from the micro structure of composites after the final processing step. In those models, a real structure of composites is substituted by the model one and analyzed to find structural parameters to fit the theoretical preconditions of the models. Detailed discussion of these approaches is far beyond the scope of this chapter and information concerning those models can be found in the original papers, or briefly described and discussed in the comprehensive review by Lux [3]. In this chapter, we will discuss in more detail only the Nielsen model. The reason is that the Nielsen model, in general, is without any doubt the most prominent model for a description and interpretation of various physical and mechanical properties of composites such as electrical conductivity, thermal conductivity as well as Young’s modulus. As noted above, the Nielsen model belongs to the group of structural-oriented models. On the other hand, taking into account classification of models of heterogeneous systems according to the different theoretical approaches, the Nielsen model belongs to the group of models based on the effective medium approximation 11.2.1.2

Nielsen’s Model

Nielsen’s model was proposed to describe an electrical and thermal conductivity of composites [36]. The most important parameters for calculating the conductivity of such composites are the aspect ratio (the lengthto-diameter ratio) of the filler and the coordination number of the filler in the mixture. The model can be expressed by the set of Eqs (11.4a,b,c): 1 + AB f σc = σm 1 − Bψ f σ f /σm − 1 B= σ f /σm + A   1 − max ψ = 1+ f 2max

(11.4a) (11.4b) (11.4c)

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Table 11.2 The values of the Einstein parameter (kE ) for common fillers. L is the length and D is the diameter of the filler. Filler

L/D

kE

spheres isotropic particles random fibers, 3D random fibers, 3D random fibers, 3D random fibers, 3D random fibers, 3D random fibers, 3D unaxially oriented fibers, parallel unaxially oriented fibers, perpendicular

1 1 2 4 6 10 15 35 – –

2.5 3.5 2.58 3.08 3.88 5.93 9.83 30.00 2L/D 0.5

where σ c , σ m and σ f are electrical conductivities of composite, matrix and filler, respectively and f is the volume fractions of filler. The constant A depends upon the shape and orientation of the dispersed particles. max is the maximum packing fraction of the dispersed particles. The most common values of parameters A and max are summarized in Tables 11.2 and 11.3. For randomly packed spherical particles A = 1.5 and max = 0.637, whereas for randomly packed aggregates of spheres or for randomly packed, irregularly shaped particles A = 3 and max = 0.637. There is a relation between the so-called Einstein coefficient kE and parameter A in the form A = kE + 1. Various values of A and m were given, for instance by Bigg [37]. Finally it must be pointed out that no percolation model reported in the literature has a universal validity. To apply any model correctly, the following parameters have to be carefully considered: (i) the size and the geometry of the conductive filler; (ii) the amount and the distribution of the filler within a matrix; (iii) interactions between filler and matrix; and (iv) the preparation method of the composites [3].

Table 11.3 Maximum packaging fractions of common filler shape. Filler

Shape

Packing

φ max

generic generic generic generic generic generic generic talc mica wollastonite asbestos glass/carbon generic generic

spheres spheres spheres spheres spheres spheres irregular flakes flakes fibers fibers fibers spheres spheres

hexagonal face centered cubic body centered cubic simple cubic random loose random close random close random 3D random 3D random 3D random 3D random 3D, L/D = 10 bimodal,* d 2 /d 1 = 7 broad

0.7405 0.7405 0.60 0.524 0.601 0.637 0.637 0.40–0.56 0.38–0.45 0.62 0.60 0.42 0.85 0.70

*Content ratio is five parts d 2 to one part d 1 .

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There are a lot of factors which complicate the correct use of the theoretical model. One of the most important is the shape of the filler. To get a solution of any model in the form of an analytical equation, the spherical shape of the filler as well as monodispersal size particles distribution is usually considered. However, it is known that a lot of frequently-used fillers, such as metal fibers, metallized glass fibers, metal flakes, graphite, carbon nanotubes or carbon black have shapes far from spherical simplification. Common shapes include rod-like, disc-like or grape-like morphology [2, 3]. It was shown through an analysis of some models that the percolation concentration strongly depends on the shape of the filler. In general, the more irregular shape of the filler leads to the lower percolation of composites [38]. For this reason, it is not surprising the fact that the lowest percolation concentrations were found in the composites filled with carbon black (fractal geometry of the chains) and with carbon nanotubes (very high aspect ratio). The effect of absolute size of the filler particles on the percolation concentration is not unambiguous, since it interferes with particle geometry and the interaction between polymer matrix and filler particles. For fibrous fillers, it was shown that an increase in the length of the filler leads to the decrease in percolation concentration [3]. However, the rigidity of the fibers have to be taken into account. Long, flexible fibers have a tendency to curl which, conversely, increases the percolation threshold. As for carbon black, an experimental result indicates that reduction of particle size leads to the decrease in the percolation threshold. An explanation of that fact is based on the statement that a decrease in the particle size leads to the increase in the total surface area of the filler and, consequently, the increase in the polymer/filler interaction result in the lowering percolation threshold [39].

11.3 Electrically Conductive Fillers 11.3.1

Carbon Black

Carbon black is an elemental carbon in a form different from diamond, cokes, charcoal and graphite. It consists of spherical-like particles linked with each other in the form of aggregates and is manufactured by the incomplete combustion of a heavy aromatic feedstock in a hot flame of (preheated) air and natural gas [40]: Cx Hy + O2 → C + CH2 + CO + H + CO2 + H2 O As can be demonstrated by using SEM, furnace carbon black occurs in the form of a continuous, oriented network, with no identifiable crystallites as shown in Figure 11.3. Such carbon black is usually not conductive and is commonly used in the rubber industry for a reinforcement of elastomers due to its high reinforcing effect. To get conductive carbon black, a heat treatment must be performed. Heat treatment causes a marked improvement in layer alignment and orientation, attaining the highly ordered capsular structure of graphitized carbon blacks, as shown in Figure 11.4. Such carbon blacks are appropriately conductive, having an electrical conductivity of the order of 104 S/m; however they are nonreinforcing [41]. Carbon black (CB) is one of the most frequently-used fillers for the improvement of electrical conductivity of polymers, despite the fact that its own conductivity is much lower than the electrical conductivity of metals or graphite (electrical conductivity of dry compressed CB is of the order of 104 S/m) [42–49]. These composites are frequently used as heaters [50], sensors for chemical vapors [51], electro-conductive rubbers [52], etc. The main advantages of CB are relatively low cost, low density and, especially, a specific structure that enables the formation of a conductive network within a matrix at low filler concentration. Two main factors influence the electrical conductivity of CB-filled composites: the content of CB in the polymer matrix and the spatial distribution of the CB particles [53]. The tendency to form network structures at low concentration

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Figure 11.3 SEM micrograph of a furnace carbon black particle without heat treatment. Reprinted from [41]. Copyright (1976) with permission from Plenum Press, New York and London.

of the filler depends on the rheology of the polymers during the mixing process and in the post-processing step, the wettability of the particles by the polymer, solidification rates after mixing and the post-processing morphology of the CB and its surface activity [53]. The basic unit of CB (termed the primary aggregate) is a permanent cluster of particles generated during the manufacturing process. Generally, aggregates do not break down during dispersion process. Agglomerates are formed when the primary aggregate particles touch the particles of a neighboring aggregate. The primary aggregates are linked each other and form a looser, more open structure. The structure can be strongly

Figure 11.4 SEM micrograph of a furnace carbon black particle after intensive heat treatment at 2700 ◦ C. The graphitized furnace carbon black shows improved layer alignment and orientation.Reprinted from [41]. Copyright (1976) with permission from Plenum Press, New York and London.

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influenced by shear forces during melt compounding with polymers. Interparticle surface forces determine the strength of agglomerates [40]. Surface area characterizes the size of the primary particles and their degree of microporosity. Hence, high surface area grades have more aggregates per unit weight, which normally results in smaller interaggregate distances and hence greater electrical conductivity at the given filler content when incorporated into the polymers. Structure may be characterized by the dibutyl phthalate absorption test, which depends on the number of particles that make up the primary aggregate and its shape [40]. So-called high structure aggregates have large, highly branched clusters of particles which also increases integrate contact, leading to higher electrical conductivity. A disadvantage of highly structured aggregates with long fibroids is that they may be susceptible to irreversible shear breakdown during blending. It is known that the carbon black may retain a significant degree of agglomeration (of size up to 10 μm) and that in highly conductive systems, the aggregates are randomly positioned in the compound rather than uniformly spaced, with a tendency to form network structures as carbon black content increased. Surface chemistry of carbon black also influences conductivity. Surface characterization is a complex field, comprehensively reviewed, for example by Bansal and Donnet [54, 55]. As a consequence of the manufacturing route, carbon blacks may be polluted by varying amounts of oxygen, hydrogen, chlorine, nitrogene, and sulfur. Even small amounts of metals (of the order of ppm) can be also detected in some grades of carbon black. The oxygen compounds can be present in a variety of structural forms, such as lactones, quinones, phenolics and carboxylics. These compounds are called ‘volatiles’ and can deteriorate the carbon black conductivity. In general, the greatest electrical conductivity of carbon black in polymers is usually found in the grades with high surface area (including high internal porosity), high structure and low volatile content [54, 55]. The different processing routes lead to the various carbon black microstructures. Image analysis, including scanning and transmission electron spectroscopy, has been widely employed to quantify aggregate dispersion and shape, which have been classified into four general categories as schematically shown in Figure 11.5.

Figure 11.5

Shape categories of carbon black aggregates.

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individual particle (a few nanometers)

agglomerate (a few tens μm)

Figure 11.6

primary aggregate

The development of carbon black morphology during mixing.

Dispergation of carbon blacks within polymeric matrices is a time-resolving process as analyzed by Medalia [56]. It was found that the conductivity of carbon black-rubber composites increases rapidly during the first stages of mixing as carbon black is incorporated and pathways are formed between the islands of rubber-filled pellet fragments. Subsequently, conductivity may decrease gradually during the later stages of mixing as the agglomerates are broken down and the gap between individual aggregates is increased. Only small changes in the relative position of neighboring aggregates can bring about a significant increase or decrease in conductivity without necessarily invoking major changes on other aspects of dispersion. This concept is schematically illustrated in Figure 11.6 which shows some amount of carbon black aggregates, initially represented as an agglomerate in the preliminary stage of dispersion, then as dispersed aggregates at an advanced stage of dispersion. The basic unit of carbon black (termed the primary aggregate) is a permanent cluster of particles generated during the manufacturing process. Generally, aggregates do not break down during the dispersion process. Interparticle surface forces determine the strength of agglomerates. At the preliminary stage of compounding of polymers with carbon black, many primary aggregates are connected to each other in the form of one big agglomerate. During mixing in the melt, the agglomerate breaks down and primary aggregates separate from each other. Finally, some primary aggregates link with each other and form a looser, more open structure. It is caused by flocculation, for instance. Agglomerates are formed when the primary aggregate particles touch the particles of neighboring aggregates. However, in reality, the contact between neighboring aggregates is not necessary for a transport of the electrons. Electrons are able to surmount a potential energy barrier that is related to the ˚ as a consequence of the tunneling effect [2, 3]. gap of the order of 15–100 A 11.3.2 11.3.2.1

Metallic Fillers Iron [57, 58]

Iron is not a common filler for an improvement of electrical conductivity of composites because of easy oxidation of iron powders; however, due to its magnetic properties, it is of interest in EMI shielding in the

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low-frequency magnetic range. More frequently, iron is used for specialty items such as magnetic tape or inks. Iron powder can be made electrolytic or by thermal decomposition of iron carbonyl. Iron powder also can be prepared by the reduction of high-purity, finely ground Swedish iron ore. Mill scale has been another important source of iron oxide for making iron powder. Other methods include the atomization of high carbon iron melted followed by decarbonization, and the atomization of molten, low carbon steel with high-pressure water jets. On the contrary to iron powder, stainless steel powder and fibers are useful in electrically conductive composites, especially the fibrous form with high aspect ratio. The ductile, yet tough, stainless steel fibers can be compounded for injection molding with minimal fiber breakage. The fine-diameter fibers are manufactured by a wire drawing process. In this process, a multiple end bundle of fibers is drawn continuously to produce a strand or tow of individual parallel fibers. The fibers start a stainless steel rod that is wrapped in alloy sheathing. The clad rods are then drawn, bundled together, and redrawn. The wire strands then undergo chemical leaching to remove the support matrix, yielding a continuous tow of multiple fibers. The process is followed by chopping of fibers to get the fibers of required length. 11.3.2.2

Copper [57, 58]

Copper and its alloys are frequently used as powders. They are prepared by several methods, including air atomization of molten copper solutions, and hydrogen reduction of copper salts in solution. The latter process requires elevated temperature and pressures. A similar process has been used for obtaining nickel and cobalt powders. The powders usually have a spherical shape; however, the grades have a special dendritic shape with high aspect ratio. Copper powders tend to oxidize and become nonconductive in polymeric composites. However, some grades are available that maintain their electrical conductivity due to special anti-corrosion surface treatment. 11.3.2.3

Aluminum [57, 58]

Aluminum powder can be prepared by grinding bulk pieces of aluminum and its alloys in a hammer mill and the particle size can be further reduced to the desired size by ball milling. The molten metal can also be atomized by high-velocity air into small spherical particles. The aluminum powder is not used to make electrically conductive plastics because the surface of the particles easily oxidizes. Fibers and flakes are used as the electrically conductive filler; however, as for copper, the filler needs special surface treatment. The fibers and flakes are often produced by a melt extraction process of solidification using a spinning wheel to throw off particles of molten aluminum alloy. The particles are immediately quenched as stable fibers of various aspect ratio. 11.3.2.4

Nickel [57, 58]

High-purity nickel powders are manufactured by carbonyl refining technology. Nickel carbonyl is converted to nickel powder: Ni(CO)4 → Ni + 4CO The powder morphology is highly specific, and different types of nickel, all ferromagnetic, are available. One type has individual spherical particles with a spiky surface with a diameter of 5 μm. Another type is characterized by strings of spherical particles with spiky surfaces, with a string containing up to 50 beads.

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The powders are often stabilized with an oxide coating and contain about 0.8% oxygen. This type of nickel is black. Nickel spheres and flakes can be coated by silver to improve their resistivity against high humidity. 11.3.2.5

Silver [57, 58]

Silver is produced in powder and flake form as a filler for the design of conductive composites. The flakes are mechanically flattened to have a very large surface and an aspect ratio. Fatty acids such as stearates and oleates are used to prevent agglomeration of the flakes. This coating can be consequently washed off to improve electrical conductivity. The flat shape of the filler is often used for a preparation of highly conductive thin film, where high electrical conductivity is caused by adjacent particles overlapping. The powder containing spherical silver particles also tends to agglomerate in common polymers (LDPE, HDPE and PS) [49, 59]. Silver is an excellent conductor and, unlike less costly metals, its oxide, sulfate, and carbonate are also good conductors [22]. In contrast, although aluminum, copper, and zinc are excellent conductors in bulk, their fine powders and flakes are susceptible to oxidation, which destroys their effectiveness because their oxides are nonconductors. 11.3.2.6

Noble Metals (Gold and Palladium)

Noble metals are seldom used as fillers in polymeric composites. It has been observed [59] that gold and palladium powders behave in considerably different ways when mixed into common polymers (LDPE, HDPE and PS). The powder containing spherical gold particles tends to agglomerate exhibiting the percolation threshold values about 11 vol. %. In contrast, the palladium powder shows threshold values above 30 vol. % probably due to very low agglomeration (Figure 11.7).

Figure 11.7 SEM micrograph of palladium particles. Reprinted from [59]. Copyright (1993) with permission from Elseiver.

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11.3.3

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Metallized Fillers

Metal-coated inorganic fibers are often used as a substitute for metals [60, 61, 62]. The advantages are the easy metallization of convenient substrates, high conductivity, low density and lower prices when compared to metals. A common application of such composites is EMI shielding materials. Among the most frequently metallized substrates are glass fibers, carbon fibers and mica [57]. Metals commonly employed for coatings are silver, copper and nickel [57]. Silver is a very convenient material for metallization, due to its high electrical conductivity, and because its oxide is also conductive and therefore the exposure to humidity does not significantly change its conductivity compared with alumina or copper [63]. Successful coatings of polymeric fillers can bring many advantages, including lower density of final filler and lower price, as well as variability in the shape of the filler [63, 64, 65]. Metallization can be carried out using an electro-less metallization method based on the oxido-reduction reaction. For example, two solutions can be used, namely one solution consisting of silver nitride dissolved in distilled water, and the other obtained by dissolving sodium-potassium tartrate in distilled water [64, 65]. The bath for electro-less deposition of silver is prepared by mixing both solutions in various volume ratios at ambient temperature. This procedure usually leads to the formation of the thin silver layer deposited on the filler surface. The thickness of the layer can be partly controlled and varies in the range from 0.5 to 1 μm. The compactness of the layer depends on the surface of the filler, especially on its polarity (the total surface free energy). It is clear that polar inorganic surfaces are coated more easily than inherently nonpolar polymeric fillers. However, an appropriate surface treatment of polymeric filler, for instance using the plasma discharge, can significantly improve polarity of those fillers and make them suitable for successful metallization. The electrical conductivity of silver coated fillers is in the range from 103 to 104 S/m. Examples of recently-developed silver-coated fillers are shown in Figures 11.8 and 11.9 [64, 65]. 11.3.4

Graphite

Graphite is a soft, crystalline form of carbon. It is one of three forms of crystalline carbon; the other two are diamond and fullerenes. Graphite is a soft mineral with a Mohs hardness of 1 to 2, and it exhibits perfect

Figure 11.8 SEM micrograph of silver-coated polyamide particles. Reprinted from [64]. Copyright (2007) with permission from Elseiver.

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Figure 11.9

439

SEM micrograph of silver-coated wollastonite fibers.

basal cleavage. Depending upon the purity, the specific density is 2.26 g/cm3 . It is gray to black in color, opaque, and has a metallic luster. It is flexible but not elastic. It has high electrical and thermal conductivity, is highly refractory, and is chemically inert [58]. There are two general types of graphite: natural and synthetic. Natural graphite occurs in metamorphic rocks such as marble, schist, and gneiss. Graphitization of naturallyoccurring organic carbon may be performed at temperatures as low as 300–500 ◦ C or as high as 800–1200 ◦ C. The three principal types of graphite–lump, crystalline flake, and amorphous – are distinguished by physical characteristics that are the result of major differences in geological origin and occurrence. Graphite exhibits the properties of metals and nonmetals, which makes it suitable for many industrial applications. The metallic properties include high electrical and thermal conductivity, whereas nonmetallic properties include chemical and environmental inertness, high thermal resistance, and lubricity. A combination of electrical conductivity and lubricity allow its use as the primary material in the manufacture of brushes for electric motors. A graphite brush effectively transfers electric current to a rotating armature while the natural lubricity of the brush minimizes fractional wear. From a thermodynamical point of view, graphite is the most stable form of carbon. This means that any other form of carbon may be transformed into graphite under suitable conditions (temperature, pressure, time). Graphite crystallizes preferably in a hexagonal lattice. Within these layers, each carbon is linked with three neighboring atoms by strong covalent forces, the distance C-C being 1.415 A. In contrast, the π link between layers is relatively weak. The distance between layers is 3.354 A [3]. Unlike monocrystals in synthetically-produced graphite, those in natural graphite have a markedly parallel orientation, which makes for a high level of anisotropy in its properties. Due to relatively weak interaction between adjacent graphitic sheets, graphite can be intercalated and finally exfoliated, similar to clays. The raw material for production of flexible graphite are large purified natural flakes with a well-ordered crystalline structure. The properties of the end product are governed mainly by the mechanical, chemical and thermal purification process used. Highly oxidizing acid mixtures, such as nitric acid and sulfuric acid, are used to produce graphite intercalation compounds. In this process, anions as HSO4 − are intercalated between the planar structures of the monocrystal [66]. These intercalation compounds can be expanded by sudden heating to high temperatures.

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Figure 11.10

SEM micrograph of natural graphite flake.

If these graphene layers could be separated down to a nanometer thickness through intercalation and exfoliation, they would form high modulus graphite nanosheets, which possess an enormous surface area (up to 2630 m2 /g) and satisfy the high aspect ratio (200–1500) criterion needed for high strength composites [67]. This process in demonstrated in Figures 11.10, 11.11 and 11.12. The natural graphite flakes are shown in Figure 11.10. After intercalation followed by thermal treatment one can get the structure as shown in Figure 11.11. Finally, in the ideal case, particles of graphite can be dispersed within a matrix as an individual sheets – Figure 11.12 [68]. Polymer/graphite nanocomposites have attracted considerable interest in both academia and industry, owing to their potential applications in advanced technologies, for example, in antistatic coatings, electrochemical

Figure 11.11

The structure of expanded graphite after intercalation followed by thermal treatment.

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Figure 11.12

441

SEM micrograph of an individual graphene sheet.

displays, sensors, catalysis, redox capacitors, electromagnetic shielding and in secondary batteries [69, 70]. Polymer nanocomposites prepared from high aspect ratio layered graphite nanofillers may achieve significant improvements in mechanical, thermal, electrical and barrier properties at very low filler concentrations [71], compared to conventional composites filled with a common type of graphite. On the other hand, the dispersion of such nanosheets in the polymer matrix is very complicated and plays a crucial role in the improvement of electrical, mechanical and thermal properties of the resultant nanocomposite [72]. The synthesis of graphite nanosheets and processing techniques that produces complete exfoliation and good dispersion of graphite particles in the matrix have been well documented in the literature. Basically, nanocomposites can be prepared by two routes: exfoliated graphite is either mixed with monomer, which then polymerizes, or blended with a polymer in the melt [73]. It was found that the presence of expanded graphite compared to regular graphite significantly decreases percolation concentration (composites containing about 2 weight % of expanded graphite was reported as electroconductive), reflected by changes in electrical conductivity, barrier properties as well as thermal stability [66]. Most approaches described in the literature are based on polymerization of monomers in the presence of expanded graphite [67]. Less attention is paid to thermoplastics filled with expanded graphite due to problematic exfoliation of graphitic sheets during blending in the molten polymer [74]. 11.3.4.1

Graphene [75]

Graphene was first obtained very recently, immediately attracting attention. Graphene, consisting of single planar layers of hexagonally arranged sp2 carbons (it can be viewed as an individual atomic plane extracted from graphite, as an unrolled single-wall carbon nanotube or as a giant flat fullerene molecule), is currently the most exciting carbon-based structure. The coupling of layers is possible due to the weak Van der Waals forces between them. Structures with up to 10 graphene layers are known as few-layer graphenes, while structures with more than 10 and less than 100 layers are considered as thin films of graphite. The number of layers is very important and needs to be controlled because it may determine the properties and performance of the graphitic material. Single graphene was obtained for the first time by Novoselov from three-dimensional graphite using

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a technique called micro-mechanical cleavage [76]. Graphene exhibits a number of exotic physical properties, previously not observed at the nanoscale. The observation of its great mechanical strength of 130 GPa, high Young’s modulus of 1 TPa, high adsorption of hydrogen and CO2 , room-temperature quantum Hall effect, ultrahigh electron mobility and ballistic transport, long electron mean free paths, zero electronic band gap, superior thermal conductivity and remarkable flexibility are among the striking properties of graphene [75]. The unusual electronic properties of this new material make it a promising candidate for future devices. While these applications are a focus of further investigations, there are some areas where graphene can be used directly. Graphene has attracted increasing attention for optoelectronic devices, super-capacitors, gas sensors, pH sensors, chemical sensors, strain sensors, biosensors, transparent films for liquid crystal devices, biodevices, DNA transistors and nanocomposite applications. The proposed applications have enhanced properties compared with carbon nanotube materials. The aforementioned method of mechanical exfoliation of graphite (also known as the ‘Scotch tape’ or peel-off method) provided a small amount of high quality samples for fundamental studies. Several other methods were subsequently developed and have been utilized to produce graphene sheets. Reported methods include the chemical reduction of liquid suspension graphene oxide (GO) [77], liquid-phase exfoliation of graphite [78], conversion of nanodiamond [79], epitaxial growth by thermal desorption of Si atoms from the SiC surface, epitaxial growth by chemical vapor deposition on transition metals, solvothermal synthesis and unzipping carbon nanotubes. Despite the number of methods for its synthesis, as-prepared graphene itself is not soluble and thus cannot be dispersed in water or any organic solvent. Suitably modified graphene nanosheets could display good solution chemistry with properties such as dispersability and solubility in water and organic solvents. Therefore, hydrophilic and organophilic affinities for graphene nanosheets should be achievable through chemical functionalization. Despite the recent progress in the production of graphene, the synthesis of uniform and large enough quantities of single-layer graphene is still an ongoing challenge. Many researchers are now focusing on derivatives of graphite, which is inexpensive and available in large quantities. A particularly popular derivative is GO, which is hydrophilic and has a larger interlayer distance than graphite. It can readily exfoliate into individual GO sheets in water and forms stable dispersions after ultrasonication. Subsequent deoxygenation via reduction can restore the electrically insulating GO to conductive graphene. Maintaining the individual separation of the graphene sheets is the most important and challenging aspect of all these synthetic routes. Bulk graphene sheets – if left unprotected – will spontaneously agglomerate and even restack to form graphite. Moreover, neat graphene is an insoluble and relatively inert material. Chemical functionalization or the use of dispersant is generally needed to prevent agglomeration. The GO synthetic pathway is attractive for stabilizing individual sheets in solutions. The oxygen functional groups that exist in GO provide reactive sites for chemical modification using known carbon surface chemistry. The chemical attachment of appropriate organic groups leads to physical separation of the resultant graphene sheets but also makes it possible to directly form stable graphene dispersions during the synthetic process. The successful dispersion of graphene has enabled the use of low cost solution processing techniques to fabricate various potentially useful graphene-based materials [80]. The GO formation involves the reaction of graphite with strong oxidizers. The introduction of oxygen containing functional groups, such as hydroxyl and epoxide, results in an increase in the interlayer spacing of GO as well as a change of hybridization of the oxidized carbon from planar sp2 to tetrahedral sp3 . GO (sometimes called graphite oxide, graphitic oxide or graphitic acid) has been known for over 150 years. It was first prepared by Brodie [81] in 1859 by repeated treatment of graphite with an oxidation mixture of potassium chlorate and fuming nitric acid. The methods most commonly used at the present time are the original Brodie synthesis, one described by Staudenmaier [82] in which the graphite is oxidized in concentrated sulfuric and nitric acids with potassium chlorate, and one described by Hummers and Offeman [83] in which the graphite is treated in a water-free mixture of concentrated sulfuric acid, sodium nitrate and potassium permanganate, respectively.

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The reduction of GO leads to graphene or materials that on the microscopic level are structurally similar to neat graphene. Different reagents (ammonia, chlorine, nascent hydrogen from aluminum powder and hydrochloric acid, hydrazine, sodium borohydride or hydroquinone) can be used and the final chemical structure of the graphene is dependent on the type of reagent. The reactions can be performed on GO nanoparticles dispersed in media, but the resulting graphene re-agglomerates. When GO is embedded in a solid matrix, the reaction leads to individually dispersed graphene nanoparticles that are unable to agglomerate (the result is a multiphase system where one phase is graphene). Electrochemical, rapid and mild thermal transformation of GO to graphene-like material was also performed. A novel approach was used in preparation of poly(arylene disulfide)/carbon nanosheet composites [75] with GO acting as an oxidizing reagent for the in situ polymerization of aromatic dithiols while the resulting reduced GO acted as the host for the composites. A similar approach was used for the production of graphene/SnO2 nanocomposites, in which GO was reduced by SnCl2 to graphene sheets in the presence of HCl and urea [75]. Similarly, the GO in the GO-titanium oxide (TiO2 ) nanocomposites was successfully reduced by UV-assisted photocatalytic reduction [75]. Graphene sheets of non-oxidized graphene, GO and reduced GO, can be further functionalized by chemical reactions with various functional groups. The carboxylic, hydroxyl and epoxy groups present on the basal plane or edges of the graphene sheets of GO allow for functionalization of the graphene by simple organic reactions. Moreover, some organic reactions can also proceed on the aromatic entities of GO or its reduced (non-oxidized) form. A simple method often used for the functionalization of graphene is based on reactions of the carboxyl groups, present in GO and located at the edges of graphene sheets, with various amines, alcohols or silanols. Reactions of the graphene carboxyl groups with amines lead to the formation of amides. For example, this type of modification was used for modification of graphene with poly(vinyl alcohol) and poly(ethyleneglycol) [75]. Some reagents, such as isocyanates, can react with both the edge carboxyl and surface hydroxyl groups of GO. The ability of isocyanates to react with hydroxyl groups, which unlike carboxyl groups are also located on the basal plane of the graphene sheets, allows for a higher degree of GO functionalization and thus better dispersability of large graphene sheets. The chemical modification of GO by aryl and alkyl isocyanates led to the hydrophobization of GO that was readily dispersible in water before modification [75]. The composites were prepared by solution phase mixing of the exfoliated phenyl isocyanate-treated GO with PS followed by their chemical reduction [75]. Such material exhibited a very low percolation threshold of fillers at 0.1 vol% as a consequence of the perfect homogeneous exfoliation and high aspect ratio of individual graphene sheets. One possible way of functionalizing graphene aromatic rings is by substitution of metallized graphite followed by electrophilic substitution of the metal graphite with halide for the preparation of exfoliated soluble graphene [75]. Another method, which does not require an oxidized graphene surface, is functionalization by diazonium salts [75]. Choi et al. [84] reported the functionalization of epitaxial graphene via nitrene radicals formed by the thermal decomposition of azidotrimethylsilane. An electrochemical approach to graphene surface functionalization was used for the preparation of ionic-liquid-functionalized graphite sheets, which could be exfoliated into functionalized graphene nanosheets [75]. A simple method for the covalent functionalization of graphene by a polymer chain is the grafting of polymer radicals onto the graphene surface such as poly(acrylic acid) and polyacrylamide [75]. In addition to all the methods of graphene surface functionalization mentioned, a hydrogenation of graphene sheets has also been reported [75]. 11.3.5

Carbon Nanotubes (CNT)

At the present time, CNT are considered as prospectively the most important filler for designing of electrically conductive composites and therefore the most intensive research effort is paid to this subject in scientific literature.

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na1

ma2

C = na1+ma2 (5,5) Armchair

Figure 11.13 The arrangement of graphene sheet and the cylinder-like structure model of the SWNT by rolling along chiral vector. Reprinted from [188]. Copyright (2009) with permission from John Wiley and Sons.

Carbon nanotubes, originally discovered in the middle of the last century [85], became the source of interest since the landmark paper by Iijima in 1991[86]. They represent attractive and challenging structures for the formation of electroconductive polymeric nanocomposites. They exhibit many interesting properties, such as extremely high mechanical strength, Young’s modulus, and strain to failure as well as remarkable electronic structure. It makes them extremely high electrically and thermally conductive [87]. Different types of CNT can be produced by various methods. The most common techniques are: arc discharge, laser ablation, chemical vapor deposition and flame synthesis. CNT consist of relatively high amount of impurities, such as catalyst residue, etc. For this reason they must be purified using the following techniques: oxidation, acid treatment, annealing, sonication and filtering [88, 89]. CNT occur in two basic distinct forms: single-walled nanotubes (SWNTs), which are composed of a graphene sheet rolled into a cylinder (Figure 11.13), and multi-walled nanotubes (MWNTs) (Figure 11.14) that consist of multiple concentric equiaxed graphene cylinders (different colors represent individual cylinders in Figure 11.14). Compared with multi-walled nanotubes, single-walled nanotubes are of greater interest owing to their expected novel electronic, mechanical, and gas adsorption properties. The electronic properties of SWCNT are the most significant characteristic of this material. SWCNT can either be metallic or

Figure 11.14 Structure of multi-walled carbon nanotubes. Reprinted from D.R. Kauffman, A. Star, Carbon Nanotube Gas and Vapor Sensors, Angew Chem Int Ed., 47, 6550. Copyright (2008) with permission from John Wiley & Sons.

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semi-conducting, depending on diameter and chirality. The chirality of nanotubes is determined by the arrangement of graphene sheets that is mathematically represented by chiral/roll-up vector (Figure 11.13). The differences in conducting character are caused by the molecular structure that results in a different band structure and thus a different band gap. The differences in conductivity can easily be derived from the graphene sheet properties [90]. Metallic and semimetallic NTs have roll-up vectors such that n − m = 3q (where q is any integer or zero) and semiconducting NTs have n − m = 3q,0. Nanocomposites based on polymers filled with CNT have a range of potential applications [91, 92, 93]. The carbon nanotubes have a unique atomic structure, very high aspect ratio, and extraordinary mechanical properties (strength and flexibility). These properties, at least theoretically, make them ideal reinforcing fibers in nanocomposites [94]. CNT-reinforced composites have been investigated for flame-retardant performance [95], improved electrical conductivity and electrostatic charging behavior, optical emitting devices [96, 97], and in lightweight, high-strength composites [95]. However, colloidal materials such as carbon nanotubes do not spontaneously suspend in polymers, thus the chemistry and physics of filler dispersion become major issues. In the case of polymers filled with carbon nanotubes, the research challenge is particularly tremendous due to the unique character of these unusual materials. Due to strong attractive interaction, nanotubes aggregate to form bundles or ‘ropes’ (Figures 11.15, 11.16) that are very difficult to disrupt [98]. In case of single nanotubes, they are only 1–3 nm in diameter; however, since they like to assemble into ropes, which consist of many nanotubes, they are most likely 10–200 nm in diameter. Furthermore, ropes are tangled with one another like spaghetti or polymers. With high shear, these ropes can be untangled, but it is extremely difficult to further disperse at the single tube level. Separated multi-walled carbon nanotubes are shown in Figure 11.17. Due to the low entropy of mixing, rigid molecules of high molecular weight require strong attractive interactions to disperse. Since the connectivity and rigidity of macromolecules drastically reduces the number of configurations available in the dispersed state, mixing becomes a problem. In the case of rigid fillers dispersed into stiff polymers, the problem is compounded in that neither species gains entropy on dispersion. It indicates that thermodynamics of CNT dispersed within a solution or within a polymer plays a significant role in an effective utilization of excellent properties of net CNT in composites. Several factors make the dispersion of CNT complicated. These factors are dominated by strong attraction between carbon species of both enthalpic and entropic origin. In addition, the low dimensionality of CNT leads to an enhancement of these attractive forces.

Figure 11.15

Bundles of aligned nanotubes – as produced.

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Figure 11.16

TEM image of partially exfoliated single-wall carbon nanotubes showing the rope-like structure.

The origin of the attractive forces between graphitic structures is well known. Due to the extended π electron system, these systems are highly polarizable, and thus subject to large attractive van der Waals forces. These forces are responsible for the secondary bonding that holds graphitic layers together. In the case of carbon nanotubes, these forces lead to so called ‘ropes’. Extended structures are formed by side-by-side aggregation of the nanotubes in ropes. When CNT are suspended in a polymer, an attractive force between fibers increases due to pure entropic factors [98]. This effective attraction is intrinsic to colloids dispersed in polymers. 11.3.5.1

Surface Modification of CNT

The notes above indicate a need for effective surface treatment of CNT. Only in such a case can they be dispersed within a polymeric matrix and significantly improve electrical, thermal or mechanical properties.

Figure 11.17

Separated multi-walled carbon nanotubes.

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In order to improve interfacial interaction between polymers and CNT, various approaches have been proposed and studied. All of these approaches are intended to modify the nature of the nanotube surface. The approaches can be categorized in general as follows: 1. Covalent surface modification of the nanotubes 2. Non-covalent surface modification of the nanotubes 3. Endohedral filling of their inner empty cavity Covalent Modification of CNT Surface The covalent modification of the nanotube surface involves chemical reactions between the nanotubes carbon atoms and the chemical reagents. These reactions cause changes in the walls of the nanotubes, and therefore cause defects in the perfect structure of the original nanotubes. Opening up of Carbon Nanotubes Through Oxidative Treatment The simplest method to open the nanotubes is their oxidative treatment. It is well known that graphite oxidizes primarily at defects of the hexagonal lattice to create etch pits. When such defect sites are present in the wall of the nanotubes, they become the center of preferential etching. However, nanotubes have additional structural features such as high curvature, helicity, and contain five and seven rings which modify the initiation and also the propagation of oxidation. Particularly for MWNT the oxidation tends to start near the tips, providing a mechanism for opening the tubes. Oxidation proceeds layer by layer, resulting in thinner tubes. Once the tip is removed, the strain induced by the distribution of pentagons is no longer there. The concentric layers of MWNT do not react at the same rate, since each shell has its own tip. It follows that the inner shells might persist longer than the outer. Therefore different oxidation rates are assumed for oxidation of an open MWNT. Upon treatment, carbon nanotubes are refluxed in concentrated HNO3 , H2 SO4 or KMnO4 solution, or at room temperature in acidic HF/BF3 solution, or even in aqueous solution of OsO4 or OsO4 -NaIO4 for 24 hours until their originally capped ends open up. The treatment in acidified KMnO4 solution seems to be a somewhat better procedure. The advantage of HF/BF3 treatment is that it is carried out at room temperature. As mentioned above, the oxidative treatment of nanotubes results in not only open nanotubes at their tips, but nanotubes thinner in diameter [99]. The extent of thinning depends on the duration of treatment. A further consequence of the oxidative treatment is the partial functionalization of the tubes, i.e. CNT become covered with carboxyl or hydroxyl groups at their ending. These functional groups make CNT partially soluble. The number or concentration of the inserted carboxyl groups can be estimated by simple acid-base titration. The concentration of the surface acid groups in the nanotubes opened by various oxidants is in the range of 2 × 1020 − 10 × 1020 site/g of nanotube [100]. The strong oxidative treatment, using HNO3 + H2 SO4 mixture for instance, can cause cutting of the nanotubes’ length. However, the carboxylic acid functional groups that are formed on the nanotubes’ surface enable them to be dispersed readily in water. Moreover, hydrophilic polymer chains can interact with these nanotubes, in order to form a composite material, where its interface is based on the hydrogen bonding between the nanotubes and the polymer matrix. The different oxidative treatment also has an effect on the final properties of polymer composites. For example, Tamburri et al. [101] observed that the conductivity of poly-1,8-diaminonaphtalene filled with SWCNT increases by increasing the CNT polarity due to the presence of –OH and –COOH groups introduced by chemical treatment in the order untreated < HNO3 < KOH < HNO3 /H2 SO4 . On the other hand, Kim and co-workers [102] observed that the conductivity of MWCNT/epoxy composites decreases as time and temperature of MWCNT oxidation in nitric acid increases. A similar result was observed for conductive poly(3-octylthiophene) [103]. The Galiotis group oxidized MWNT with reagents of different oxidation power [104], and they investigated mechanical and electrical properties of MWNT/epoxy composites with respect to the chemical oxidation of the nanofiller [105]. The stronger oxidative treatment applied to the nanotubes (the oxidation in boiling HNO3 ) brings about a two-fold

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increase of the flexural modulus and a decrease of DC conductivity by two orders of magnitude. Combined improvement of both modulus and DC conductivity can only be achieved by mild oxidative treatments such as that provided by the NH4 OH/H2 O2 reagent. It is also clear that the presence of carboxyl groups on the CNT surface open new possibilities to perform chemistry on these terminal groups. Several applications, for instance the use of CNT as polymer fillers, require or take advantage of the presence of chemical groups situated on their surface or at their tip. Functionalization by Organic Compounds There are many ways showing how to modify polarity of CNT utilizing previous carboxylation of CNT. In many cases, low polarity of CNT is required to disperse them in unpolar polymeric matrices. In this case, long alkyl chains covalently linked onto the CNT surface can help to improve quality of CNT/polymer dispersion. One possibility of how to realize it is as follows: a reaction of the carboxylated CNT with thionyl chloride and followed by a reaction with an amine terminated molecule leads to higher hydrophobicity of CNT, dependently on the length of alkyl chain (R). However, R can represent various organic functional groups. The nature of the functional group, and in particular the R functional group in the amine reagent, will determine the dispersion properties of the modified nanotubes. Similar reactions can provide alcohols, hydroxyl terminated polymers (polyols or polyesters), and isocyanate terminated polyurethanes [106]. The oxidized nanotubes can be also used for attachment of initiator for polymerization and followed atom transfer radical polymerization of acrylates and styrenes [106] or reversible addition-fragmentation chain transfer polymerization of styrenes and acrylates [106]. Other chemical modifications of the nanotubes’ surface can be achieved by various approaches. One example is the plasma treatment of aligned and separated nanotubes, to cause functional groups to be attached covalent to the surface. Another example is the reaction of an aryl radical on the nanotubes’ surface [107], hydrogenation, cycloaddition of carbenes, nucleophilic addition, ozonolysis or mechanochemical functionalization (the ball-milling of MWNT in reactive atmospheres was shown to produce short tubes containing different chemical functional groups such as amines, amide, thiols and mercaptans [108]). The polymer-bound carbon nanotubes can be formed by covalently attaching nanotubes to highly-soluble linear polymers, such as poly(propionylethylenimine-co-ethylenimine) (PPEI-EI) via amide linkages or poly(vinyl acetate-co-vinyl alcohol) (PVA-VA) via ester linkages. The samples of polymer-bound nanotubes are soluble in both organic solvents and water, and highly colored homogenous solutions are formed [109, 110]. It has been demonstrated that carbon nanotubes can be covalently fluorinated within the temperature range from 250 ◦ C to 400 ◦ C [111]. After that, fluorine can be effectively removed from the SWNTs using anhydrous hydrazine. Boul et al. reported a sidewall-alkylated nanotubes modification employing a sidewall-fluorinated nanotubes reaction with alkyl magnesium synthesis or by reaction with alkyllithium precursors [112]. The product of reaction between alkyllithium and nanotubes can serve as initiator for anionic polymerization of styrene [106] or methacrylate [106]. For applications requiring the high conductivity of carbon nanotubes, all the above methods are not very attractive due to losing CNT conductivity. For this reason, the non-covalent procedures were proposed. Non-Covalent Modification of CNT Surface This approach to the surface modification of nanotubes is based on the van der Waals attraction between the nanotubes and various molecules. The advantage of this procedure is that it does not break CNT during treatment as well as not disturbing delocalized π electrons, and thus it does not change an inherent electrical conductivity of CNT. A common example is a derivative of pyrene which can be attached non-covalently to the sidewalls of nanotubes. The molecule consisting pyrenyl group can interact with the nanotube sidewalls, and on the other side the molecule has a functional group (the succinimidyl ester group, for example) that can attract amine functional groups of the polymer matrix [113].

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Another mechanism of non-covalent modification of the nanotubes’ surface is called wrapping. In this case, specific polymer chains are wrapped around the nanotubes in a helical shape. In the wrapping mechanism, the interactions between the nanotubes and the polymer matrix in the composite material can be controlled, by choosing and modifying the functional groups along the wrapped polymer chains. For instance, Stoddart et al. [114] produced bundles of SWNTs that have a conjugated polymer helically wrapped around. The polymer is a poly(m-phenylenevinylene) with octyloxy chains. After adding SWNT to a solution of this polymer, the mixture was sonicated resulting in a formation of a stable suspension of nanotubes. The polymer attaches to the nanotube surfaces, as a result of π – π interactions [114]. In this section we have devoted some time to the surface modification of CNT. The reason for this is not only that CNT are the most intensively studied fillers for an improvement of electrical conductivity of polymeric composites, but also those approaches can be applied for a surface treatment of any carbon fillers, such as carbon black or graphite. Endohedral Filing of Carbon Nanotubes with Polymers The encapsulation of PS into hollow CNTs was reported by Liu [115]. In this approach, monomer and radical initiator were first carried into the cavities of the CNTs with the aid of supercritical CO2 . Steinmetz [116] used a supercritical fluid impregnation when MWCNTs were filled with a photo-conducting PVK and a conducting polymer polypyrrole prepared inside the nanotubes by free radical polymerization of N-vinyl carbazole and oxidation polymerization of pyrrole, respectively. A similar procedure was used for preparation of MWCNTs filled by polyacetylene [117]. Bazilevsky described a new open-air method for filling MWCNTs by diffusion at room temperature [118]. Relatively low molecular weight polymers, such as PEO (600 kDa) and PCL (80 kDa) were encapsulated in as-grown wettable MWCNTs with 50–100 nm diameters. As expected, relatively small flexible polymer molecules entered the nanotubes while large macromolecules (∼1000 kDa) remained outside. 11.3.5.2

Composites Based on CNT

A myriad of various polymeric composites blended with CNT have been prepared, studied and reported in the scientific literature and there is not space to review it in this chapter. Very briefly, polymer/CNT composites can be divided into three main groups: (i) composites prepared in solution by in situ polymerization of monomer/CNT mixtures; (ii) composites prepared by mixing CNT and polymeric solution; and (iii) composites prepared from thermoplastic polymers blended with CNT in molten state. It was demonstrated that in situ polymerization is the best way for preparation of homogenous dispersions because CNT are more likely to disperse in a precursor monomer than in the polymer which usually leads to the lower percolation concentration of CNT within a matrix [119–122, 123]. On the other hand, this approach is not industrially attractive for large-scale applications. For this reason, a big effort is devoted to effective dispersion of CNT within common thermoplastic matrices such as polyethylene, polypropylene, polyamide, etc, during blending in common mixing devices [124–127]. 11.3.6

Conducting Polymers

Conducting polymers are sometimes called ‘synthetic metals’ because they possess electric, electronic, magnetic and optical properties as metals or semiconductors, while partly preserving the mechanical properties of conventional polymers. Conducting polymers were pioneered by Alan J. Heeger; Alan G. MacDiarmid and Hideki Shirakawa discovered polyacetylene, modified by oxidation with chlorine, bromine or iodine vapor [128]. Many interesting applications were found in the following years, such as light-emitting diodes, solar cells, transistors, diodes, holographic storage media, chemical and biological sensors, capacitors, batteries,

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Poly-trans-acetylene NH NH

NH

NH

NH

NH

Polypyrrol S

S S

S

S

S

Polythiophene NH

NH

NH

NH

Polyaniline

Poly(p-phenylene)

Figure 11.18

The chemical structure of the most common electrically conducting polymers.

anti-static coatings, electromagnetic shielding, anti-corrosive coatings, gas and liquid separation membranes, artificial muscles, lithography and metallization, photo-electrochromic devices and xerographic photoreceptors [129, 130, 131]. Among the most frequently-studied conducting polymers besides polyacetylene belong polyaniline, polythiophene (and various polythiophene derivatives) and polypyrrole [132, 133]. Whereas polyacetylene is sensitive to atmospheric oxygen and to humidity, polyaniline, polypyrrole and polythiophene may be synthesized directly in the doped form and they are very stable in air. In intrinsically conductive polymers, conductivity is attributed to delocalization of π -bonded electrons, in conjugated double bonds, over the polymeric backbone. These materials also exhibit unusual electronic properties, such as low energy optical transitions, low ionization potentials and high electron affinities [134]. Figure 11.18 shows the structure of some conducting polymers. In order to become electrically conductive, the polymer has to be disturbed either by removing electrons from (oxidation) or by inserting electrons into (reduction) the material. This process is known as doping. Depending on the dopant, a conducting polymer exhibits either p-type or n-type conductivity. For example, treatment with halogens has been termed ‘doping’ by analogy with the doping of semiconductors. During this process, an organic polymer, either an insulator or semiconductor having a low conductivity, typically in the range 10−10 to 10−5 S cm−1 is converted into a polymer which is in the ‘metallic’ conductive regime (∼10 S cm−1 ). Dopant ions are generally introduced into the polymer system during chemical or electrochemical polymerization and they play an important role in balancing the charge distribution within the polymer. The dopant ion influences the distribution of positive charge on the polymer backbone and, when constrained over a small area, this can cause additional closely-spaced electronic levels to form within the band gap. Small lattice distortions occur locally so that conformational changes as well as charge inequalities (defects of bond conjugation) are induced in the polymer backbone. The net effect is that the oxidation state changes and the equivalent of free radicals can be formed. A charged site interacting with a free radical forms a polaron. Polarons are radical cations or anions that are generated during the doping process. Especially in

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polymers with non-degenerated ground state, the polarons are highly unstable but they can be stabilized by further oxidation of the polymer to form radical dications – bipolarons [135]. These conjugational defects lie within the band gap of the material and it is these defects that make conducting polymers interesting for chemical sensing. The introduction of dopants and the creation of bipolarons influence the electrical conduction mechanism. The transport of charge most often consists in phonon-assisted tunneling (hopping) of charge carriers between localized states induced by conjugational defects. Increasing density of the localized states leads to intra-chain carrier delocalization (along the polymer chain) accompanied by growing charge carrier mobility and later the delocalization is extended into a 3D inter-chain electron transfer. Instead of the doping level, the conductivity in conducting polymers (density of localized states) can be controlled also by modifications of polymer backbone, e.g., using copolymeration of the conducting polymer with its derivatives [136].

11.3.7

Fillers Coated by Conducting Polymers

Despite the fact that conducting polymers may also be used as the filler into the nonconductive polymers, it is better to use them for coating of inorganic or organic particles or fibers. For example, polypyrrole was recently investigated as a prospective agent for electroconductive coating of different inorganic as well as organic fibers. Much attention has been paid to a treatment of carbon and graphite fibers [137, 138, 139], glass fibers [140], as well as partly to polymeric fabrics and fillers [141, 142]. Polypyrrole was also employed for a treatment of cellulose-based materials. Bjorklund and Lundstrom described the synthesis of paper/PPy composites [143] with a conductivity of about 2 Scm−1 . Miˇcuˇs´ık et al. [144] prepared highly pure cellulose fibers coated with polypyrrole (PPy) and consequently used them for preparation of the composites based on the LDPE, HDPE and PCL matrices. The recent production of a new conducting composite material based on polyaniline and wood by simple one-stage synthesis from aniline and wood sawdust was described [145] as well as composites of low density polyethylene (LDPE) filled with canadian switch grass coated with polypyrrole were prepared and studied [146]. The SEM micrograph cellulose fibres coated with 20 wt.% of polypyrrole is shown in Figure 11.19.

Figure 11.19 SEM micrograph of cellulose fibers coated with polypyrrole. Reprinted from [144]. Copyright (2006) with permission from John Wiley & Sons.

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11.4 Effect of Processing Conditions on the Electrical Behavior of Composites Polymeric composites are most often prepared in the molten state using convenient mixing equipment. It is known that the blending is a complex process which has a strong influence on the final properties of composites. Filler particles incorporated into polymer melt significantly affect the melt flow characteristics, including viscosity, melt strength and a die swell. These effects are more pronounced at the highest filler content. A detailed description of these effects is discussed, for example in the work of Hornsby [147]. Very generally one can say that in the case of electrically conductive composites their conductivity as well as a percolation concentration strongly depends on the processing conditions. Later we will discuss very briefly some features of a preparation of electrically conductive composites in the melt. 11.4.1

Blending of Polymeric Composites

Two basic mechanisms of mixing in the melt is distinguished, namely extensive and intensive (sometimes called dispersive) mixing. In the case of mixing very viscous melts, extensive mixing is achieved by convection, which can be distributive or laminar and which leads to the obtaining of uniform dispersion. Distributive mixing causes a rearrangement of the components of mixture through an ordered or random process, such as in the pre-blending of components in the solid state. Laminar mixing, however, is achieved by imposing on the materials permanent deformation in various laminar flow patterns such as shearing, squeezing or elongational flow. This will necessitate mixing in the melt state through the imposition of large strains and is generally accompanied by an increase in interfacial area between components in the system. Effective laminar mixing is strongly influenced by the initial orientation and spatial location of solid components [148, 149] Intensive mixing normally involves a rupture of agglomerates formed by a solid state. Within a polymer melt, dispersive mixing of a minor particulate phase (filler) is effected by shear stresses at the polymer melt/filler interface. Breakdown of the structure is accompanied by distribution of the separated particles throughout the polymeric matrix through and intensive mixing step. Dispersive mixing of particulate filled polymers is therefore influenced by a variety of factors relating to machine design and operation, together with material composition [150]. The mixing process of fillers into a polymer involves the following stages: Filler wetting: Polymers wet the filters and penetrate into their void spaces; this results in a replacement of air bubbles within a filler agglomerates by polymeric phase. The polymer is believed to penetrate not only the void space of the agglomerate, i.e. between the aggregates, but also within the aggregates. It is particularly important for highly structurized carbon black. It was proposed that in elastomers filled with carbon black rubber which fills the void space within each aggregates occluded and immobilized, thus acting as part of the filler rather than a part of the deformable matrix [151]. Dispersion: This effect relates to the break-up of the agglomerates and separation of the resulting fragments to a stage where re-agglomeration will not occur. It is generally accepted that agglomerates will break when internal stresses, induced by viscous pressure on the particles, exceed a certain threshold value [152]. When these forces exceed this critical value which is equal to the attractive cohesive forces, the aggregates break apart. A few models were developed to describe this process, e.g. [153]. Distribution: Once the filler agglomerates are broken, separation of the closely spaced agglomerate fragments and their distribution throughout the polymeric matrix is accomplished. Accordingly, the degree of the filler dispersion no longer depends on the stress, but only on the total level of strain applied to the matrix

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(for example, laminar hear strain). On the other hand, it has also been suggested that even with well separated aggregates of the filler (particularly carbon black), flocculation may occur by diffusion in the hot polymeric system contributing to the formation of a network of particles and electrically conductive structure. The notes above indicate that the mixing procedure has a strong influence on the electrical conductivity of composites. Obviously, the extent of processing conditions on the final conductivity strongly depends also on the morphology of the filler. It is particularly important for carbon black, where it was found that the conductivity of the composites filled with highly porous carbon black were much less sensitive to the processing time than the conductivity of composites filled with structural carbon black due to its tendency to structural breakdown. The type of polymer also has a strong influence on the conductivity of composites. It was found that the greater degree of internal phase separation developed during polymer solidification leads to the higher conductivity of the resulting composites. For example, polypropylene, having a spherulitic morphology of crystallites has a strong influence on the segregation of carbon black particles. It leads to the lowering of percolation concentration. Similarly, a phase-separated amorphous copolymers or terpolymers such as ABS have an ability to form segregated networks.

11.4.2

Effects of the Secondary Processing Steps on Conductivity

It is generally accepted that post-processing steps also have very significant influence on the final electrical conductivity of polymeric composites. For instance, a high shear secondary processing step such as injection molding can result in dramatically lower electrical conductivity relative to low shear stress formed during compression molding. It is caused by the filler orientation during injection. Higher shear forces evocate higher orientation of the filler which results in the increase in composite resistivity. This effect is more pronounced in the skin than in the core [154]. The influence of shear forces on the orientation of polymeric chains and on the orientation of network of the filler strongly depends on the molecular weight of polymer. At low and intermediate molecular weights the conductivity decreases after injection molding and material orientation. It has been shown that the decrease in conductivity in extruded samples against the compression molded samples is caused by an increase of orientation of the anisotropic carbon black aggregates. However, as the polymer molecular weight increases, the shear and elongational stress contribution in the matrix during injection molding increases, leading to a breakdown of the aggregates and, as a consequence, reduced anisotropy. This has the effect of decreasing conductivity in isotropic compression molded due to a less defined percolation conductive filler network. However, for injection molded materials, particularly using the highest molecular weight polymer, even though the filler is highly destroyed segregation (meaning carbon black clusters or aggregates), flow induced segregation of those filler particles becomes dominant, generating highly conductive channels, and a value of conductivity is significantly higher than that for unoriented materials with a nearly homogeneous particle distribution. Annealing is another important effect having an influence on the structure of the conductive network. This process, mainly under pressure, may lead to the reformation of the network and to the decrease of composite resistivity. Finally, in practical preparation of electrically conductive composites in a common mixing equipment various processing parameters, such as shear rates, a number and temperature of mixing zones, screw profile, degree of fill and others have to be taken into account. A detailed description of this issue can be found in related literature [155, 156].

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Effect of Polymer Characteristics on the Electrical Conductivity of Composites

The selection of polymeric matrix depends on the final application of electrically conductive composites. Functional applications of the basic polymer is defined overall by complex mechanical, optical, thermal, ecological and other properties that are determined by the material in its condition of use. The operating temperature and a presence of adverse environmental factors are also taken into account. Many polymers are negatively influenced by high temperature, mechanical stress or irradiation dose. Some plastics absorb water from the atmosphere during use and this influences the electrical and mechanical properties. Various oils, acids and other substances badly affect some plastics. At dynamic loads, the loading rate, temperature, size and form of the articles as well as a development of microdefects become critical. From this reason, the selection of the appropriate polymer must be considered very carefully for any concrete application. Among physical and structural parameters that influence the final electrical behavior of materials are, mainly, the surface chemistry of polymers, degree of crystallinity and molecular weight. 11.4.4

Crystallinity Effect

It is clear that when filler is dispersed within a polymeric matrix, particles can be distributed only in the amorphous part of the matrix. It means that the crystalline portion has a strong influence on the percolation concentration of the filler within a matrix, as demonstrated in Figure 11.20. Amorphous PS, LDPE with degree of crystallinity of 38 wt.% and HDPE with degree of crystallinity of 75 wt.% were blended with the same carbon black grade (Vulcan XC-72). As can be seen in Figure 11.20, the percolation concentration decreases with an increase in degree of crystallinity [157]. The same tendency is observed for electrical conductivity of composites. It is seen that an increasing crystallinity of the matrix results in an increase in electrical conductivity of investigated composites at comparable filler content.

14 12 10 8

logσc/σm

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PS/CB composites, Φc=0.09 LDPE/CB composites, Φc=0.06 HDPE/CB composites, Φc=0.02

2 0 -2 0,00

0,05

0,10

0,15

0,20

0,25

volume portion of CB

Figure 11.20 The dependence of electrical conductivity of composites (σ c ) filled with carbon black (CB). σ m is the electrical conductivity of neat matrix.

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On the other hand, the presence of the filler particles within a matrix may have an influence on the change of degree of crystallinity as well as on the size of crystallites. When carbon black was blended with polypropylene, it was shown [158] that the microcrystalline structure was not significantly affected by the filler. The size, but not the geometry, of the spherulites was affected. Isothermal crystallization studies using DSC are used to investigate the influence of the filler on the degree of crystallinity and on the size of crystallites. On the one hand, the filler can act as an nucleating agent which results in an increase of the crystalline portion. It is detected as an increase of the specific enthalpy of melting. On the other hand, the filler particles can reduce the degree of crystallinity due to steric hindrances. The particles may suppress the free movement of polymeric chains, which leads to the restriction of folding ability. It results in a decrease of degree of crystallinity, detected by the decrease in a specific enthalpy of melting. As well as the change of degree of crystallinity, particles can also influence the size of crystallites. The filler particles usually decrease the size of crystallites which is detected as the decrease of melting temperature [159]. 11.4.5

Effect of Polymer–Filler Interaction

The interfacial compatibility between polymer and filler is closely associated with the wettability and quality of dispersion of the filler. Due to their completely different chemical and structural characters, polymers have much lower polarity than inorganic electrically conductive fillers, such as metals and metal oxides, as well as carbon graphitic structures such as graphite, graphitized carbon black and carbon nanotubes. Whereas a total surface free energy of polymers is of the order of a few tens mJ/m2 , a total surface energy of common inorganic is of the order of a few hundreds mJ/m2 . For instance, a total surface free energy of carbon black varies from 70 mJ/m2 to several hundred mJ/m2 , dependent on the composition and processing conditions. As for polymers, a total surface free energy of most polymers has a value in the range 30–40 mJ/m2 [160]. Large differences in surface tension between polymers and fillers result in less compatibility between components as well as in worse quality of dispersion. For this reason, to improve compatibility and dispersability of the filler, surface treatment of fillers, polymers or both is frequently performed. Surface modification of filler includes either a simple treatment of surfaces by surfactants such as fatty acids, or various covalent or non-covalent surface modifications using various low molecular compounds or even special polymers. The surface treatment of carbon fillers (especially carbon nanotubes) will be discussed in the next section in more detail. Here we will briefly review some polymer modifications. 11.4.5.1

An Improvement of Polymer/Filler Compatibility

Modification Through Adding of Low Molecular Compound Organic acids, particularly fatty acids, are extensively used as surface treatment systems for particulate inorganic fillers [161]. Without the treatment, the filler particles would shrink excessively, forming strongly bound agglomerates that are difficult to disperse in the matrix. Fatty acid modification causes the filler surface to become hydrophobic and moisture adsorption during storage is significantly reduced. Fatty acids are also effective in assisting incorporation of polar mineral fillers in nonpolar polymer matrix melts. This results in reduced melt viscosity and improved dispersion, in most situations. The type of surface treatment used also affects the final product properties. As well as improved dispersion, reduced melt viscosity results in lower shear degradation. There are two main methods of surface treatment addition. The filler may be pre-treated, where the surface treatment is pre-adsorbed onto the filler surface before incorporation into the matrix can be added directly to the pre-mix of filler and matrix, prior to melt blending. The most known fatty acid for surface treatment is stearic acid and its salts, e.g. with Zn or Ca [161]. It improves dispersability of inorganic filler within a polymeric matrix. In this case, the chemical bond between filler and polymer is not formed. It means that fatty acids are non-interacting surface modifiers. They reduce adhesion between filler and matrix. In highly filled composites this reduction in adhesion

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usually leads to an increase in toughness accompanied by a decrease in ultimate strength. On the other hand, the strength can be positively influenced due to reduction of probability of aggregation of particles. This is important mainly for very small particles, which have a significant tendency to form aggregates. Usually, 0.2–2 weight % of it in final composite is commonly used [162]. It can be added to the blend directly or in the batch. Modification Through Grafting of Polymers with Specific Functional Groups Maleated polymers are among the widest known family of functionalized polymers used as compatibilizer and adhesion promoters [163]. Anhydride groups can react with amine, epoxy and hydroxyl groups. Epoxydized polymers are also commercially available products. Generally they are mainly modified by glycidyl methacrylate. They are very reactive with NH2 , anhydride, acid, hydroxy groups. They are recommended to compatibilize polyesters (PET, PBT) and olefinic polymers or elastomer [164]. Nonreactive polar copolymers (specific interaction by polarity) are also employed to reduce interfacial tension and increase the adhesion by creation of a specific polar interaction like hydrogen bonding or Van der Waals forces. The compatibilizer has to be compatible with one phase (generally nonpolar) and has to create specific interactions with the other phase [165]. Ionomers also have a strong potential for an improvement of polymer/filler compatibility. Ionomers are polymers containing attached ionic groups to the backbone chain. They are copolymers, containing both nonionic repeat units, and a small amount of ion-containing repeat units. The ionic groups make up less than 15% of the polymer. The possibility of introducing different ionic groups (either acid or neutralized form) to the backbone chains enables to create a vast number of different structures that can be employed for various compatibilizations of immiscible polymer–polymer systems as well as polymer–inorganic filler systems due to specific electrostatic interaction. Ionic groups have significant influence on the polarity of polymeric backbone and therefore improve mutual adhesion between polymers and fillers. One typical example of an ionomer is poly(ethyleneco-methacrylic acid). This polymer is a sodium or zinc salt (which provides the ions) of copolymers derived from ethylene and methacrylic acid. Moreover, they could be employed as coupling agents during mixing of components in extruder during on-step processing, similarly yo maleated polymers, e.g. [166]. 11.4.6

Multiphase Morphology of Polymers and Its Influence on the Conductivity of Composites: Multipercolation Effect

The distribution of the filler within a polymeric matrix can be very effectively influenced by the use of multiphase polymer blends. Through a careful selection of the polymer type in the blend, together with their relative concentrations and miscibilities, it is possible to produce a variety of microstructures, such as immiscible droplets phase morphologies or co-continuous distribution of the phases [167]. An introduction of the filler (all the known papers deal with carbon black) into one of the polymer phases or into the phase interface has an dramatic influence on the electrical conductivity of composites as well significantly lowering the percolation concentration. Lipatov [168] investigated conductive polyethylene/polyoxymethylene blends prepared by dilution of a composite polyoxymethylene/carbon black in the second polymeric component. This resulted in localization of the filler at the interface between the polymeric components. It was found that the carbon black was largely expelled to the interface of the polyoxymethylene phase as it crystallized, although some remained dispersed in this component. The consequence of this conductive network at the polyethylene/polyoxymethylene interface was an increase in conductivity level at lower filler content. Gueskens [169] found that the electrical conductivity of natural rubber/polyethylene blends filled with carbon black was much higher than that for the individual components at the same loading level. The same

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author reported later the results on polystyrene blended with ethylene-propylene random copolymer mixed with carbon black. It was found that these blends were much more conductive than each of the polymer constituent at the same filler content. It was also shown that the addition of carbon black to these two immiscible polymers reduced the size of the dispersed phase. Sumita et al. [170, 171] studied electrical conductivity of three types of blend system filled with carbon black, namely HDPE/PP, PP/PMA and HDPE/PMMA blends. Two types of carbon black distribution were observed: one where the filler is homogeneously incorporated in one phase of the blend, and the other where the filler is preferably localized at the interface between the two polymers. The electrical conductivity was determined both by the filler concentration in the filler-rich phase, and also the structural continuity of this phase, i.e. so-called double percolation. This approach leads to the formation of highly conductive polymeric composites with significantly reduced percolation concentration. Another route for influencing electrical conductivity in thermoplastics containing carbon black has been investigated using core-shell graft copolymers [172]. The composite was formed from elastomeric cores of crosslinked polyethylene-vinyl acetate copolymer (EVA) and crosslinked polybutylacrylate (BA) with a polyvinylchloride (PVC) and filled with carbon black. It was found that less than half the amount of carbon black was needed in the core-shell copolymer system to achieve a given level of conductivity, than with conventionally formulated PVC-based composites without a loss of flexibility of the final material.

11.5 Applications [2, 3, 57] 11.5.1

EMI

A requirement for the control of EMI emanating from computers, radios, cell phones and other equipment significantly stimulated development of electrically conductive polymeric composites. Electromagnetic radiation is absorbed and reflected by highly conductive materials, such as metals. For many reasons, plastics replaced metals in housing of computers and other electronic devices. However, radiation passes through plastics so has a negative influence on the electronics as well as on the surroundings. To overcome this problem, the plastic articles were coated with conductive materials by metal flame spray, vacuum metallization and conductive paints. The conductive paints are formulated with resins such as acrylics, epoxy or urethane filled with conductive fillers. The shielding composites have the value of electrical conductivity of the order of 1 to 2 /sq. It results in a shielding efficiency of 40 to 60 dB. The surface resistivity of shielding composites is usually in the range from 0.1 to 100 /sq. Generally, the higher this value, the better the EMI protection afforded. Most applications require a minimum of 30 dB attenuation, which would prevent 99.9% of the incoming radiation from passing through the shielding. Today, most articles needing EMI shielding protection are designed using common thermoplastics and thermosetting resins blended with aluminum, stainless steel, metalized glass and graphite. 11.5.2

ESD

ESD (electrostatic discharge) is another growing area of electroconductive composites, since electrostatic discharge can destroy electronic components and also cause fire or explosion in operating rooms or around flammable solvents. An electrical charge is accumulated as a result of the rubbing or sliding of dielectric materials. Under some conditions, mainly at low humidity, the induced charge can build up to 30,000 V, which can discharge as an arc or spark when a person comes in contact with it. The conductivity of composites for ESD application may be lower than that for EMI shielding. Its value is in the range from 102 to 107 /sq. The antistatic composites fall in the 109 to 1013 /sq range.

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Electrically Conductive Adhesives

Electrically conductive adhesives are formulated by mixing thermosetting resins with good adhesion to metals or other required substrates and highly conductive filler. Just for an illustration, fine silver powder has been applied in many commercial adhesives where high conductivity is required. U.S. Patent 2,444,034 (1948) describes a conductive adhesive using silver powder mixed with a solution of cellulose acetate or nitrate. Another composition, patented in U.S. Patent 2,774,747 (1956), was formulated from epoxy resin, silver flakes, ultrafine precipitated silver powder and diacetone alcohol. The epoxy-based adhesives are made as one component or two-part systems (resin and hardener). An importance issue with all the adhesives is to achieve a balance between adhesive strength (as measured by ASTM lap shear method, for instance), volume resistivity and a filler content. As for other fillers used in the design of electrically conductive adhesives, gold, nickel, copper and metallized glass fillers were frequently employed. Another type of adhesive is called hot-melt adhesive. In this case, thermoplastics having a good adhesion to the required substrates are used as the matrix. A good adhesion to metals or to inorganic substrates can be achieved by grafting those thermoplastics with appropriate functional groups such as maleic anhydride, ionomer-based groups, etc.

11.5.4

Conductive Rubbers

Conductive rubbers are frequently used in automobile door handles, bumpers, fenders, air bag covers; wire & cable, films, gaskets, hoses, pipes, thermal tapes and conveyor belts. Various types of carbon black are almost exclusively used as the filler.

11.5.5

Semi-Conductive Cable Compounds

Compounds with volume resistivity of 1–100 ohm.cm are used in regular cable compounds. HDPE is usually used as the matrix; frequently crosslinked by peroxides or X-ray whereas CB is used as the filler.

11.5.6

Fuel Cells

A fuel cell is an electrochemical device that combines hydrogen and oxygen to generate electricity with water and heat as its by-product. The polar plates used in certain types of fuel cells need to be electrically conductive. Different kinds of graphite are usually used as the filler whereas epoxy or phenolic resins are used as matrix.

11.6 Resistance Measurements The conductivity of these electrically conductive composites ranges over many orders of magnitude. It is obvious that the measurements of electrical properties must be done accurately. For that reason, it is important to pay attention to the proper arrangement of the experiments. One experimental setup may be suitable for samples with high conductivity, while another for samples with low conductivity. Of course, a better and deeper understanding of the conducting polymer nature demands not only the proper implementation of the experiments, but also a good interpretation of measured data. The aim of this section is to provide the reader with an overview of basic experimental techniques of the DC measurements that enable information to be obtained about the correct conductivity of the material under investigation.

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Rp

Rc Rsp

Figure 11.21 The total resistance includes resistance of the probe Rp , resistance of the contact Rc and spreading resistance Rsp .

11.6.1

Two-Probes Method

First of all, some comments on two-probe measurements will be made. In this method, each probe has two roles, serving as a voltage as well as a current probe. If we analyze the contact originated in the probe attached to the surface of a measured sample, the resistance Rp of the probe, the resistance Rc of the contact, and the spreading resistance Rsp must be taken into account (Figure 11.21). One can see that total resistance RT between these two probes can be written as: RT = U/I = 2R p + 2Rc + 2Rsp + Rs

(11.5)

where Rs is the true resistance of the material, U is the potential difference between the probes, and I is the passing current. The multiplication factor of two originates from the number of contacts (Figure 11.22). If both probes are identical in shape, in contacting force, and in the material they are made from, one can assume both Rc as well as Rsp values to be identical for both probes. The contact resistance Rc originates from the contact between the probe at the surface of the measured material, the spreading resistance Rsp is related to the flow of the charge carriers from one probe into the material and from this material into the second probe. The probe resistance Rp can be determined separately (two probes in the set-up can be short-cut and then their resistance can be determined). The determination of the resistances Rc as well as Rsp will be discussed later.

I+

I– U

Rs

Figure 11.22

The two-probe configuration.

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Figure 11.23

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The four-probe configuration.

Four-Probes Method

Two-probe measurement techniques, where two electrodes are used simultaneously as voltage and current probes, cannot provide the possibility of determining a correct value of Rs in Eq. (11.5). The solution of this problem is based on the separation of current and voltage probes, i.e. another set of two probes must be added to the experimental set-up (Figure 11.23). Outer probes are usually used as current probes; inner probes are similarly used as voltage ones. In this case, the current-carrying probes do not influence the determination of true resistance of the material, because the voltage drop is measured on voltage probes localized at different places. The measurements of voltage are very often done by employing a voltmeter with very high input impedance to minimize the drawing current (e.g. using an electrometer). In such a way, no probe resistance, spreading resistance or contact resistance would influence the measurements. Therefore, a method in which the current and voltage probes are kept apart is involved! For material scientists it is much more interesting to describe these materials by the resistivity instead of a resistance, the former parameter being an intensive property independent of the sample dimensions and geometry. The resistivity then allows them to compare the electrical properties of various materials. To determine the resistivity from the known resistance, a geometrical factor of the measured sample has to be taken into account. If the shape of the sample is a bar-like type with a uniform cross-section along the length of the sample (Figure 11.24), the resistivity can be calculated as: ρ=

ab R z

(11.6)

where the resistance is R = U/I. Homogeneity of the material is also assumed. As a rule, the voltage U is measured in volts [V], the current I in amperes [A], and the resistance R in ohms [ ], the sample dimensions a, b and c are given in centimeters [cm], and then the resistivity is obtained in ohm.cm units [ cm]. Let us add that the conditions for correct measurement on a bar-like sample should fulfill the conditions: z ≈ c – 2a, z ≈ 4a. This is necessary to avoid the nonuniformity of the current profile entering the sample. On the other hand, the distance z has to be long enough to get a measurable value of the voltage. A longer distance between voltage probes implies a higher value of voltage. The contact at the place, where the current enters the sample, should be ohmic and, therefore, a large area contact is obviously needed. Another way of determining the resistivity of the sample is the method introduced in semiconductor studies in the 50s by Valdes [173]. This method is suitable for samples with shapes other than the bar-like geometry

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c

z

b U a I

Figure 11.24

Measurement set-up for bar-shaped sample.

mentioned above. A so-called standard four-probe method is based on the determination of the potential drop under probe touching the measured material in an ideal point contact, where current enters the sample (Figure 11.25): dU =

ρI dr 2πr 2

(11.7)

Outer probes carry the current to the sample again – flow-in and flow-out currents are equal in magnitude but opposite in direction (Figure 11.26); inner probes record the difference U between potentials U + and U − :

U

+

s1 =− s1 +s2

U

−

s 2 +s3

=−

ρI ρI dr = 2 2πr 2π ρI ρI dr = 2πr 2 2π

s3

Hence, U = U + − U − =

ρI 2π







1 1 − s1 s1 + s2 1 1 − s2 + s3 s3





1 1 1 1 + − − s1 s3 s1 + s2 s2 + s3



I

dr r

Figure 11.25

The potential drop on increment dr caused by current I.

(11.8)

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U

I+

s1

Figure 11.26

I–

s2

s3

Outer probes carrying current and inner probes serve as potential probes.

For the resistivity ρ we can easily calculate: ρ = 2π

U I



1 1 1 1 + − − s1 s3 s1 + s2 s2 + s3

−1 (11.9)

For s1 = s2 = s3 ≡ s we get: ρ = 2π s

U I

(11.10)

This equation is exactly valid for a semi-infinite sample only and this is not the case in practice. When the distances between the sample periphery and the probe spacing are comparable, correction functions must be involved as a multiplicative factor to the right side of Eq. (11.10). This will be commented on later. Also the ratio of the sample thickness to the probe spacing s has to be considered. As a limit case for the thickness t of the sample, we can discuss two examples. The first is for the sample with t  s; this case is described by Eq. (11.10). The second is for an opposite case, t  s; then for the thin-sheet form of the sample, the difference U between potentials U + and U – has a form:   (s1 + s2 )(s2 + s3 ) ρI ln U = U − U = 2π t s1 s3 +

−

(11.11)

and for the probe spacing s1 = s2 = s3 ≡ s, we get a resistivity independent of s: ρ=

π t U ln 2 I

(11.12)

Please note that ratio ρ/t, where t is sample thickness is called the sheet resistance and defined as the resistance between opposite edges of a square with an arbitrary size. The relationship given by Eq. (11.12) can be rewritten as: Rsq =

π U ∼ U = 4.532 ln 2 I I

(11.13)

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I–

I+

+

–

463

2t

2t I

U=0

I s

t

s

s

jn = 0 2t

I–

I+

I+

I–

2t

Figure 11.27 The images for determination of the correction function. (The condition of the conducting boundary is depicted.)

This quantity is expressed in the units of ohms per square [ /square]. The measurements on thin layers become important as new materials for modern technology go to ever thinner films. Very thin layers need to be supported by other substrates, because alone they have poor mechanical integrity. The supports can influence the resistivity of a thin layer under investigation. Thus, two cases have to be considered: (i) the resistivity of the substrate is considerably higher than the resistivity of the deposited thin layer, or (ii) its resistivity is considerably lower. In these situations, corrections have to be done to avoid mistakes during thin layer resistivity determination. Relevant correction functions can be derived by using calculations offered by Valdes [173, 174]. They used the procedure known as ‘method of images’ (Figure 11.27), where an infinite number of imaginary sources is introduced on both sides of the thin layer. The accuracy of the correction function F(t/s), where t/s is the ratio of the layer thickness and probe distance, depends on the number of imaginary sources taken into the calculation: 

 −1/2  2 −1/2 −1 ∞ s2 s 4s n 2 2 (−1) + 4n − + 4n Fcon (t/s) = 1 + 4 t n=1 t2 t2

(11.14)

Fcon is correction function for the case of the conducting substrate, for an isolating substrate. The correction function Fiso can be derived in the form: 

 −1/2  2 −1/2 −1 ∞ s2 s 4s Fiso (t/s) = 1 + 4 + 4n 2 − + 4n 2 t n=1 t2 t2

(11.15)

Both functions are displayed in Figure 11.28. The value of the correction function for a certain ratio t/s represents a multiplication factor in the right side of Eq. (11.10). Two comments should be added. (1) For t/s > 4, both correction functions are nearly the same within 1% difference, and (2) for t/s < 0.5, the values of the function Fcon become very high due to the electrical short-cut in the conducting substrate.

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102 Fcon

Fcon, Fiso

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100

Fiso 10–1 0.1

Figure 11.28

1 t/s

10

The correction functions according to Eqs (11.14) and (11.15).

The situation is more complicated when the dimensions of measured samples are comparable with the probe spacing s. Let us consider the case of in-line arrangement of the probes, which are located close to the straight sample periphery. We must take into account two configurations. One is the perpendicular direction to the nonconductive sample periphery; the other is parallel to the sample periphery. The shape of periphery correction function Fnor for a normal direction is:

Fnor

    −1     2k −1 2k −1 2k −1 2K −1 = 1+ 1+ − 2+ − 4+ + 5+ s s s S

(11.16)

where k is the distance of the outer probe to the sample periphery. For a parallel direction the shape of periphery correction function Fpar is:

F par =

⎧ ⎨ ⎩

1+2 1+



2k s

2 −1/2

⎫  2 −1/2 ⎬−1 k − 1+ ⎭ s

(11.17)

where k is now the distance of all probes to the sample periphery. The correction functions, as well as probe configurations, are shown in Figure 11.29. From the practical point of view, it is worth emphasizing that the error introduced to Eq. (11.10) is about 1% for k/s > 3. For the case of a conducting periphery, the sign of image terms (Figure 11.29) has to be changed to get the true value of correction functions for both probes’ directions, a normal:

 Fnor

    −1     2k −1 2k −1 2k −1 2k −1 = 1− 1+ + 2+ + 4+ − 5+ s s s s

(11.16a)

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Figure 11.29 The correction functions according to Eqs. (11.16), (11.16a), (11.17) and (11.17a). Inserts – image method for the conducting boundary (lower curves in main figure).

as well as a parallel one:

 = F par

⎧ ⎨ ⎩

1−2 1+



2k s

2 −1/2

⎫  2 −1/2 ⎬−1 k + 1+ ⎭ s

(11.17a)

We have described so far two relative positions of the probes and sample periphery. Of course, the real number of such positions is much greater; details can be found in the literature. The corrections also depend on the shape of the sample. The configuration of the probes need not be linear. A square configuration was also considered, and then the new correction has to be determined [175]. 11.6.3

Van der Pauw Method

In the methods described above, the shape of the sample must be known. The mutual position of probes is also important for the accurate determination of the true value of resistivity. Sometimes it is not possible to prepare a sample with a defined shape. In such cases, using conformal mapping, Van der Pauw solved the problem of the resistivity determination on the sample of an arbitrary shape [176]. The conditions which must be satisfied are: (1) the flat sample is without isolated holes (from a mathematical point of view the sample is singly connected); (2) the sample fulfils the condition ρ/t = constant, where t is sample thickness; and (3) four point contacts are located in any location on the periphery of the measured sample. The resistance RABCD is defined as RABCD =

UCD IAB

(11.18)

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Figure 11.30 The sample of arbitrary shape with contacts for resistivity measurements according to the van der Pauw method.

where I AB denotes the current flowing through contacts A and B, U CD is voltage drop measured between the contacts C and D (Figure 11.30). When the positions of current and voltage contacts are shifted, the resistance RBCDA is defined in a similar way. Then, the resistivity is given by: ρ=

π t RABCD + RBCDA ln 2 2

(11.19)

Equation (11.19) holds only for a symmetrical sample such as a circle or a square with symmetrically located contacts. In other cases, Eq. (11.19) must be multiplied by implicit function FvdP defined by:  cosh

Q − 1 ln 2 Q + 1 FvdP



1 = exp 2



ln 2 FvdP

 (11.20)

where Q = RABCD /RBCDA . The function FvdP depends only on the ratio Q and can be calculated numerically (see Figure 11.31). The best way to make the measurements in the van der Pauw arrangement is to use four springing knife-like contacts. Sometimes, another type of contact is used, e.g. alloyed ones. It can be hardly considered to be point

1.0

0.8

Fvdp

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0.4

0.2 1

10

100

1000

Q

Figure 11.31

The correction function for the van der Pauw method.

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C

φ

D d3 φ

C

B

B

Figure 11.32

D

d2

A

467

A

C

A

B

Real contacts placed out of their proper position.

contacts. Also, the position of contacts cannot be exactly on the periphery of the sample. That is why some corrections for non-point contacts, as well as their incorrect position, have to be taken into account. Van der Pauw derived the changes of resistivity for the line contacts on the periphery of the circular sample as [154]: d12 ρ ≈− ρ 162 ln 2

(11.21)

where  is the sample diameter, d1 is contact elongation along the periphery (Figure 11.32 (left)). For the contact elongation in normal direction to the periphery of the sample, the changes of the resistivity can be written as d2 ρ ≈ − 22 ρ 4 ln 2

(11.22)

where d2 is contact elongation perpendicular to the periphery (Figure 11.32 (middle)). If the contact is in the distance d3 from the sample edge, then the change in the resistivity is d2 ρ ≈ − 23 ρ 2 ln 2

(11.23)

(see Figure 11.32 (right)). Rigorous studies of a variety of sample shapes used for the van der Pauw contact configuration have been done. For routine material characterization, the clover-leaf sample is usually preferred (Figure 11.33 (left)). This shape provides a large area for contact preparation, but a special procedure is needed for this unusual shape adjustment. For the monitoring of the material resistivity during integrated circuit technology, simpler shapes are required, e.g. squares (Figure 11.33 (middle)) or crosses (Figure 11.33 (right)). It has been shown [177] that, for a square-shaped sample, the positioning of contacts in the middle of the side is better than at the corners. The Greek cross is also popular because of its simple shape and small influence of the contact size as well as the contact position. For the sample structure represented in Figure 11.33 (right) with L = W, r ≤ L/6, where r is the distance of the contact from the shoulder edge of the cross, the resistance can be easily determined using the relationship [178]: ρ=

πt RABCD ln 2

(11.24)

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L

D D

A

C

W

A r

C

B B

Figure 11.33

Various sample shapes used for the van der Pauw method.

where RABCD is defined in Eq. (11.18), and t is sample thickness. When the thickness is unknown, we can determine the ratio ρ/t that is called a sheet resistance (13). 11.6.4

Spreading Resistance of the Contacts

In the contact analysis of Eq. (11.5), the spreading resistance Rsp has been mentioned. This resistance is associated with a current crowding in the vicinity of the contact region of the material under the attached probe as illustrated in Figure 11.21. When one probe has a large area contact with a small resistance placed on one side of the sample, the other contact is needle-like on the opposite side; the current flowing through the material is constrained in the vicinity of the needle-like probe. Closer to large area contact, the current lines spread out. As a consequence, the resistivity of the material ρ determines the value Rsp Rsp =

ρ 2d

(11.25)

where d is the diameter of the nonpenetrating cylindrical ohmic contact attached on a plane semi-infinite sample surface. Such description is idealized and should be modified by a multiplier k slightly depending on ρ [179]. In practice, k must be determined by the calibration process using standard samples. Also a barrier contact resistance Rb can influence the measured value and should be taken into account where contact is not perfectly ohmic. Thus, finally, the measured value can be written in the form [180]: Rmeas = Rb +

ρ k(ρ) 4a

(11.26)

To accurately determine the spreading resistance Rsp the load of the probe must be taken into account as well as mechanical properties of the probe. Electrical and mechanical properties of the measured material also affect the resulting value. For a hemispherical probe with a tip radius a, Eq. (11.25) must be replaced by Rsp =

ρ 2πa

(11.27)

The concept of spreading resistance is widely used, particularly for the profiling of inorganic semiconductors, namely for silicon. First of all, the choice of the probe material must be done with respect to its

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bevel edge

I

469

–

I

beveled surface

original surface S s

substrate

Figure 11.34

The two-equal-probes experimental set-up for spreading resistivity measurements.

mechanical and electrical properties. The next step is represented by sample-surface preparation. Furthermore, the mechanics of the contact preparation between the probe and a sample, as well as the problem of reproducible probe moving should be solved. In a short time, two equal probes placed on the same side of the sample have replaced a simpler experimental set-up consisting of the single moveable probe together with a large-area ohmic contact on the back of the sample. For two-probe arrangement 2Rsp is measured instead of Rsp , of course. The practical realization of such feasible experimental set-up for profiling is depicted on Figure 11.34. Two movable probes that are raised up and lowered measure the resistance; as a result, we can get the doping profile in a sample as a function of the depth. Note that the beveled angle is typically about 1 ◦ . When minimal step of probes moving is within 10 μm, the spatial resolution of profiling is less than 200 nm! To determine the concentration profile from measured spreading resistance, the correct model of contacts as well as a suitable algorithm must be available. The voltage between the probes during measurement is usually less than 25 mV (i.e. less than kT/e), about 5 mV typically. The current-voltage characteristics of the contact barrier are therefore nearly linear. In such selection of voltage, the contact resistance Rc is reduced and the spreading resistance Rsp is dominant in Eq. (11.5). The advantage of such a small voltage is also in low Joule-heating of the region under the probes. The important role of accurate determination of the probe radius a is evident from Eqs (11.25)–(11.27), but it is not easy to determine the true value. As a possible procedure the relationship [181] can be used Rsp =

s  ρ ln πt a

(11.28)

where t is sample thickness. The probes must be well separated, i.e. several probe spacing from the bevel edge. Knowing the resistivity ρ, the spreading resistance Rsp is measured several times with different probe spacing s before the correct probe-radius a can be determined. The contact resistance (Eq. (11.5)) is realized after attaching the probe to the surface of the material under investigation. The contact between the probe and a material has either linear or nearly linear current–voltage characteristics or nonlinear characteristics. The first is called an ohmic contact; the second is called a rectifying contact. As a rule, the probe is made from the metallic material, and the rectifying contact is named as a Schottky contact. The crucial importance of the rectifying behavior of the contact is a value of the work function of the probe material, p , as well as the work function of the material under investigation, m . For p /m <1, the contact is ohmic, for p /m > 1, the contact is rectifying. For rectifying contact barrier

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d

t

Figure 11.35 The sample with the point contact of circular shape on the upper side and the large area contact on the bottom-side.

height, φ B appears, the value of the barrier height being defined as φB =

m −χ e

(11.29)

where e is electron charge, and χ is the electron affinity of material. The sample with a point contact of a circular shape from one side and a large-area ohmic contact on the back side can be analyzed according to the method described in [182] and successfully applied by other authors [183, 184]. There, the total resistance RT between the point and large-area contacts is summarized as RT = Rc + Rsp + Rs + Rbc

(11.30)

where the resistance of the back contact Rbc is usually small compared to other terms in Eq. (11.30) and, therefore, Rbc can be neglected. The resistance of the material Rs depends on material thickness t and resistivity ρ, and both parameters are known. Resistance Rs is also often neglected. The spreading resistance Rsp in contact arrangement illustrated in Figure 11.35 is done by [182] Rsp =

ρ B 2d

(11.31)

where the correction factor B is approximated by   2 4t B ≈ arctan π d

(11.32)

where d is the contact diameter. For d/t → 0, B approaches the unity; for d/t > 10, B goes to (8/π )× (d/t)–1 ∼ = 2.5465×(d/t)–1 as can be seen in Figure 11.36. 11.6.5

Contact Resistance

The contact resistance Rc is the ratio of the specific contact resistance ρ c and the contact area; for a circular area contact Rc =

4ρc π d2

(11.33)

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B

1

0,1

0,01 0,1

1

10

100

d/t

Figure 11.36

The correction function according to Eq. (11.31).

The total resistance RT in Eq. (11.5) is a quotient of the potential drop between the small circular-point contact and a large-area contact divided by the current passing through the sample. This quantity RT is measured. Back contact resistance Rbc as well as resistance of the material Rs are neglected. Then, RT – Rsp is a linear function of the reciprocal contact area 4/π d2 . The specific contact resistance ρ c is then the slope of this linear function. When the sample thickness t as well as the sample resistivity ρ are known, ρ c can be determined by using several circular contacts with varying diameters. Another contact arrangement can be used according to Figure 11.37, i.e. both contacts are located on one side of the sample. In such case, the total resistance RT is RT = 2Rc + 2Rsp + Rs

(11.34)

where the individual terms have a meaning described above. Detailed analysis has been done for the structure depicted in Figure 11.38, where the cross-section and top view are shown. Two main differences have to be mentioned: the contact shape is rectangular and the sample thickness t is very small. In such a case, the total resistance RT can be written as RT =

I+

t

Figure 11.37

Rsq b + Rd + Rw + 2Rc W

I–

d

d

The sample with two lateral contacts from one side.

(11.35)

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I–

t

W

Z

b

Figure 11.38

A thin sample with two rectangular contacts in cross-section and top view.

where Rsq is the sheet resistance, b is the distance between contacts, W is the sample width, Rd is the resistance due to current crowding under the contacts, Rw is the resistance due to contact width correction when Z < W, last term Rc is the contact resistance. The resistance due to current crowding Rd is estimated as [185]: Rd =

2Rsq t ln 4 W

(11.36)

When the contact with Z is different from the sample with W and condition Z  W − Z is fulfilled then Rw is Rw =

Rsq π



1 k+1 k2 − 1 ln + ln k k−1 k2

 (11.37)

where k = W/(W – Z). Then, Rc can be determined when RT is measured and other terms in Eq. (11.30) are calculated using the known Rsq of the material. This method is continually spread out. The multiple-contact method has been involved as well as the transfer-length method but the characterization of these methods is out of the scope of this chapter. Also, contact-less methods can provide information about material resistivity. A great advantage for such measurements consists in no destruction of some parts of the measured sample during contact formation. Contact-less methods are based on the interaction of the electromagnetic radiation with free charge-carriers in materials. Electromagnetic radiation is used in the different frequencies, e.g., radio waves (frequency about 107 Hz), microwaves (frequency about 1010 Hz), and infrared radiation (frequency about 1014 Hz). Such broad ranges of frequencies need different experimental set-ups. Also some difficulties connected with the necessity of the calibration process must be overcome. For more information, the reader should search the specialist literature, e.g. [186, 187].

References 1. A.K. Jonscher, Dielectric Relaxation in Solids, Chelsea Dielectric Press, London, 1983. 2. C.C. Ku and R. Liepins, Electrical Properties of Polymers, Chemical Properties, Hanser Publishers, 1987. 3. F.J. Lux, Mater. Sci., 28, 285–301, 1993.

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4. V.E. Guˇl, Structure and Properties of Conducting Polymer Composites, VSB BV, Utrecht, 1996. 5. P. Bujard, K. Munk and G. Kuehnlein, Thermal conductivity of a chain of particles in close contact in a matrix of epoxy resin, in Thermal Conductivity 22, T.W. Tong (Ed.) Technomic Publishing Co. Inc., Lancaster, 1994. 6. G. Wu, J. Lin, Q. Zheng and M. Zhang, Polymer, 47(7), 2442–2447. 2006. 7. W. Thongruang, R.J. Spontak and C.M. Balik, Polymer, 43(8), 2279–2286. 2002. 8. I. Nov´ak, I. Krupa and I. Chod´ak, Eur. Polym. J., 39(3), 585–592, 2003. 9. J.M. Hammersley, Proc. Cambridge Phil. Soc., 53, 642, 1957. 10. S.R. Broadbent and J.M. Hammersley, Proc. Cambridge Philos. Soc. 53, 629–645, 1957. 11. D. Stauffer and A. Aharony, Introduction to Percolation Theory, Taylor and Francis, London, 1992. 12. M. E. Fisher J. Math. Phys. 2, 620–627, 1961. 13. S. Kirkpatrick, Rew. Mod. Phys., 45, 574–588, 1973. 14. R. Zallen, The Physics of Amorphous Solids, John Wiley & Sons Inc., NY, 1983. 15. B. Derrida, D. Stauffer, H.J. Herrmann and J. Vannimenus, J. Physique Lett., 44, 701–706, 1983. 16. I. Balberg and N. Binenbaum, Phys. Rev. B, 35, 8749–8752, 1987 17. A. Malliaris and D.T. Turner, J. Appl. Phys., 42, 614–618, 1971. 18. S.F. Lascelles and S.P. Armes, J. Mater. Chem., 7, 1339–1347, 1997. 19. J. Prokeˇs, I. Kˇrivka and J. Stejskal, Polym. Int., 43, 117–125, 1997. 20. D.S. McLachlan, J. Appl. Phys., 70, 3681–3682, 1991. 21. L. D’Ilario and A. Martinelli, J. Mater. Sci. Lett., 10, 1465–1467, 1991. 22. R. Kuˇzel, I. Kˇrivka, J. Kub´at, J. Prokeˇs, S. Neˇsp˚urek and C. Klason, Synth. Met., 67, 255–261, 1994. 23. J. Janzen, J. Appl. Phys., 46, 966–969, 1975. 24. R. Schuler, J. Petermann, K. Schulte and H.P. Wentzel, Macromol. Symp., 104, 261–268, 1996. 25. M. Sumita, A. Asai, N. Miyadera, E. Jojima and K. Miyasaki, K, Coll. Polym. Sci., 264, 212–217, 1986. 26. K. Miyasaka, K. Watanabe, E. Jojima, H. Aida, M. Sumita and J. Ishikawa, J. Mater. Sci., 17, 1610–1616, 1982. 27. B. Wessling, Die Makromo.l Chem., 185, 1265–1271, 1984. 28. B. Wessling and H. Volk, Synth Met, 15, 183–193, 1986. 29. T. Slupkowski, Phys. Status. Solid, 38, 329–333, 1984. 30. C. Rajagopal and M. Satyam, J. Appl. Phys., 49, 5536–5542, 1978. 31. S.K. Bhattacharya and A.C.D. Chaklader, Polym. Plast. Techno.l Engn., 19, 21–51, (1982). 32. I. Kˇrivka, J. Prokeˇs, E. Tobolkov´a and J. Stejskal, J. Mater. Chem., 9, 2425–2428, 1999. 33. K. Yoshida, J. Phys. Soc. Jpn., 59, 4087–4093, 1990. 34. R.L. McCullough, Comp. Sci. Technol., 22, 3–21, 1985. 35. G. Ondracek, Metall, 36, 523–531, 1982. 36. L.E. Nielsen, Ind. Engng. Chem. Fund., 13, 17–20, 1974. 37. D.B. Bigg, Adv. Polym. Sci., 119, 1–30, 1995. 38. G.E. Pike and C.H. Seager, Phys. Rew., B10, 1435–1446, 1974. 39. A.I. Medalia Carbon Black-polymer Composites: The Physics of Electrically Conducting Composites, pp. 1–49. Sichel E.K. (Ed.), Marcel Dekker Inc., New York, 1982. 40. J.B. Donnet and A. Voet, Carbon Black: Physics, Chemistry and Elastomer Reinforcement, Marcel Dekker Inc., New York, 1976. 41. J.A. Manson and L.H. Sperling, Polymer Blends and Composites, Plenum Press, NY, 1976. 42. F. Gubbels, S. Blacher, E. Vanlathem, R. Jerome, R. Deltour, F. Brouers and Ph. Theyssie, Macromolecules, 28, 1559–1566, 1995. 43. M. Kozlowski, Polym. Networks Blends, 4(4), 163–172, 1995. 44. M. Narkis and A. Vaxman, J. Appl. Polym. Sci., 29, 1639–1652, 1984. 45. Y. Bin, C. Xu, D. Zhu and M. Matsuo, Carbon, 40, 195–199, 2002. 46. L. Flandin, A. Hiltner and E. Baer, Polymer, 42, 827–838, 2001. 47. M. Omastov´a, S. Podhradsk´a, J. Prokeˇs, I. Janigov´a and J. Stejskal, Polym. Degrad. Stab., 82(2), 251–256, 2003. 48. M. Hinderman-Bischoff and F. Ehrburger-Dolle, Carbon, 39, 375–382, 2001. 49. J. Kub´at, R. Kuˇzel., I. Kˇrivka, P. Bengtsson, J. Prokeˇs and O. Stefan O, Synth. Met. 54, 187–194, 1993. 50. F. El-Tantawy, K. Kamada and H. Ohnabe, Mat. Lett, 56(1–2), 112–126, 2002.

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79. 80. 81. 82. 83. 84. 85. 86. 87. 88. 89. 90. 91.

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Handbook of Multiphase Polymer Systems E.S. Tillman and N.S. Lewis, Sensors Actuat: B-Chem., 96(1–2), 329–342, 2003. N.C. Das, T.K. Chaki and D. Khastgir, Carbon, 40(6), 807–816, 2002. D. Pantea, H. Darmstadt, S. Kaliaguine and Ch. Roy, Appl. Surf. Sci., 217(1–4), 181–193, 2003. R.C. Bansal, J.B. Donnet and F. Stoeckli, Active Carbon, Marcel Dekker, Inc., New York, 1988. J.B. Donnet, R.C. Bansal and M.J. Wang, Carbon Black Science and Technology, Marcel Dekker, Inc., New York, 1993. A.I. Medalia, Rubber Chem. Technol., 59, 432 454, 1986. H.S. Katz and J.V. Milewski, Handbook of Fillers, 1st Edition, Springer, London, 1987. G. Wypych, Handbook of Fillers, 2nd Edition, William Andrew, Toronto, 1998. R. Kuˇzel, J. Kub´at, I. Kˇrivka, J. Prokeˇs, C. Klason and O. Stefan, Mat. Sci. Engn., B17, 190–195, 1993. H. Dai H.H. Li and F. Wang, Surf. Coat. Tech., 201(6), 2859–2866, 2006. I.J. Youngs, N. Bowler, K.P. Lymer and J. Mussain, J. Phys. D.: Appl. Phys., 38(2), 188–201, 2005. B. Agoudjil, L. Ibos. J.C. Majest´e, Y. Candau and Ye.P. Mamunya, Compos. Part A-Appl S, 39, 342–351, 2008. H.C. Kim and T.L. Alford, J. Appl. Phys., 94(8), 5393–5395, 2003. I. Krupa, G. Mikov´a, I. Nov´ak, I. Janigov´a, Z. N´ogellov´a, F. Lednick´y and J. Prokeˇs, Eur. Polym. J., 43, 2401–2413, 2007. V. Cecen, A. Boudenne, L. Ibos, I. Nov´ak, Z. N´ogellov´a, J. Prokeˇs and I. Krupa, Eur. Polym. J., 44, 3827–3834, 2008. A. Herold, D. Petitjean, G. Furdin and M. Klatt, Mat. Sci. Forum, 152, 281–288, 1994. A. Celzard, J.F. Mareche, G. Furdin, Prog. Mater. Sci., 50, 93–179, 2005. SGL Carbon datasheet. A. Malinauskas, Polymer, 42, 3959–3972, 2001. T.A. Skotheim, R.L. Elsenbaumer and J.R. Reynolds, Handbook of Conducting Polymers, Marcel Dekker, New York, 1998. M. Yeng, In: Polymer Nanocomposites, Y.-W. Mai, Z.-Z. Yu (Eds), CRC Press, New York, 2006. A. Yasmin, J.J. Luo and I.M. Daniel, Compos. Sci. Technol., 66, 1182–1189, 2006. G. Chen, C. Wu, W. Weng, D. Wu and W. Yan, Polymer, 44, 1781–1784, 2003. X.Y. Chen, J.W. Shen and W.Y. Huang, J. Mater. Sci. Let.t, 21, 213–214, 2002. ˇ Z. Spitalsk´ y, M. Danko and J. Mosn´acˇ ek, Curr. Org. Chem., 15, 1133–1150, 2011. K.S. Novoselov, A.K. Geim, S.V. Morozov, D. Jiang, Y. Zhang, S.V. Dubonos, I.V. Grigorieva and A.A. Firsov, Science, 306, 666–669, 2004. S. Stankovich, D.A. Dikin, R.D. Piner, K.A. Kohlhaas, A. Kleinhammes, Y. Jia, Y. Wu, S.T. Nguyen and R.S. Ruoff, Carbon, 45(7), 1558–1565, 2007. Y. Hernandez, V. Nicolosi, M. Lotya, F.M. Blighe, Z.Y. Sun, S. De, I.T. McGovern, B. Holland, M. Byrne, Y.K. Gun’ko, J.J. Boland, P. Niraj, G. Duesberg, S. Krishnamurthy, R. Goodhue, J. Hutchison, V. Scardaci, A.C. Ferrari and J.N. Coleman, Nat. Nanotechnol., 3(9), 563–568, 2008. O.E. Andersson, B.L.V. Prasad, H. Sato, T. Enoki, Y. Hishiyama, Y. Kaburagi, M. Yoshikawa and S. Bandow, Phys. Rev. B, 58, 16387–16395, 1998. D. Li and R.B. Kaner, Science, 320, 1170–1171, 2008. B.C. Brodie, Phil. Trans., 149, 249–259, 1859. L. Staudenmaier, Ber., 31(2), 1481–1487, 1898. W.S. Hummers and R.E. Offeman, J. Am. Chem. Soc. 80(6), 1339–1339, 1958. J. Choi, K.-J. Kim, B. Kim, H. Lee and S. Kim, J. Phys. Chem. C, 113 (22), 9433–9435, 2009. L.V. Radushkevich and V.M. Lukyanovich, Zur. Fis. Chim., 26, 88–95, 1952. S. Iijima, Nature, 354, 56–58, 1991. T.W. Ebbesen, Carbon Nanotubes: Preparation and Properties, CRC Press, Boca Raton, FL, 1997. S.Iijima, C. Brabec, A. Maiti and J. Bernhole, J., J. Chem. Phys., 104, 2089–2092, 1996. E. Farkas, M.E. Anderson, Z.H. Chen and A.G. Rinzler, AG, Chem. Phys. Lett., 363(1–2), 111–116, 2002. V.N. Popov, Mat. Sci. Eng.: R: Reports, 43(3), 61–102, 2004. A.G. Rinzler, J.H. Hafner, P. Nikolaev, L. Lou, S.G. Kim, D. Tomanek, P. Nordander, D.T. Cobert and R.E. Smalley, Science, 269, 1550–1553, 1995.

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475

W.A. de Heer, A. Chatelain and D. Ugarte, Science, 270, 1179–1180, 1995. P.G. Collins, A. Zettl, H. Bando, A. Thess and R.E. Smalley, Science, 278, 100–103, 1997. H.D. Wagner, O. Lourie, Y. Feldman and R. Tenne, Appl. Phys. Lett., 72, 188–190, 1998. T. Dagani, Chem. Eng. News., 23, 11–13, 1999. P.D. Calvert, Nature, 399, 210–211, 1999. S.A. Carren, P.M. Ajavan, W.J. Blau, D.L. Carroll, J.N. Coleman, A.B. Dalton, A.P. Davey, A. Drury, B. McCarthy, S. Maier and A. Strevens, Adv. Mater., 10, 1091–1093, 1998. C. Bower, A. Kleinhammes, Y. Wu and O. Zhou, Chem. Phys. Lett, 288, 481–486, 1998. N. Yao, V. Lordi, S.X.C. Ma, E. Dujardin, A. Krishnan, M.M.J. Treacy and T.W. Ebbesen, J. Mater. Res. 13, 2432–2437, 1998. B.C. Satishkumar, A. Govindaraj, J. Mofokeng, G.N. Subbanna and C.N.R. Rao, J. Phys. B: At. Mol. Opt. Phys. 29, 4925–4934, 1996. E. Tamburri, S. Orlanducci, M.L. Terranova, F. Valentini, G. Palleschi, A. Curulli, F. Brunetti, D. Passeri, A. Alippi and M. Rossi, Carbon, 43(6), 1213–1221, 2005. Y.J. Kim, T.S. Shin, H.D. Choi, J.H. Kwon, Y.Ch. Chung and H.G. Yoon, Carbon, 43(1), 23–30, 2005. E. Kymakis and G.A.J. Amaratunga, J. Appl. Phys., 99, 08432, 2006. V. Datsyuk, M. Kalyva, K. Papagelis, J. Parthenios, D. Tasis, A. Siokou, I. Kallitsis and C. Galiotis, Carbon, 46, 833–840, 2008. ˇ Z. Spitalsk´ y, C.A. Krontiras, S.N. Georga and C. Galiotis, Composites A, 40, 778–783, 2009. ˇ Z. Spitalsk´ y, D. Tasis, K. Papagelis and C. Galiotis, Prog. Polym. Sci. 35, 357–401, 2010. J.L. Bahr, J. Yang, D.V. Kosynkin, M.J. Bronikowski, R.E. Smalley and J.M. Tour, J. Am. Chem. Soc, 123, 6536–6542, 2001. D. Tasis, N. Tagmatarchis, A. Bianco and M. Prato, Chem. Rev., 106, 1105–1136, 2006. S.S. Wong, E. Joselevich, A.T. Woolley, C.L. Cheung and C.M. Lieber, Nature, 394(2), 52–55, 1998. J.E. Riggs, Z. Guo, D.L. Carroll and Y.P. Sun, J. Am. Chem. Soc., 122, 5879–5880, 2000. J.E. Riggs, D.B. Wakker, D.L. Carroll and Y.P. Sun, J. Phys. Chem. B., 104, 7071–7076, 2000. O. Lourie and H.D. Wagner, Comp.Sci.Tech., 59, 975–977, 1999. R.J. Chen, Y. Zhang, D. Wang, and H. Dai, J. Am. Chem. Soc., 123, 3838–3839, 2001. A. Star, J.F. Stoddart, D. Steuerman, M. Diehl, A. Boukai, E.W. Wong, Y. Yang, S. Chung, H. Choi and J.R. Heath, Angew. Chem. Int. Ed., 40, 1721–1725, 2001. Z. Liu, X. Dai, J. Xu, B. Han, J. Zhang, Y. Wang, Y. Huang and G. Yang, Carbon, 42(2), 458–460, 2004. J. Steinmetz, S. Kwon, H.J. Lee, E. Abou-Hamad, R. Almairac, C. Goze-Bac, H. Kim and Y.-W. Park, Chem. Phys. Lett., 431(1-3), 139–144, 2006. J. Steinmetz, H.J. Lee, S. Kwon, D.S. Lee, C. Goze-Bac, E. Abou-Hamad, H. Kim and Y.-W. Park, Curr. Appl. Phys., 7(1), 39–41, 2007. A.V. Bazilevsky, K. Sun, A.L. Yarin and C.M. Megaridis, Langmuir, 23, 7451–7455, 2007. J. Sandler, M.S.P. Shaffer, T. Prasse, W. Bauhofer, K. Schulte and A.H. Windle, Polymer 40, 5967–5971, 1999. X.Y. Gong, J. Liu, S. Baskaran, R.D. Voise and J.S. Young, Chem. Mater., 12, 1049–1052, 2000. M.S.P. Shaffer and A.H. Windle, Adv. Mater., 11(11), 937–941, 1999. Z.J. Jia, Z.Y. Wang, C. Xu, J. Liang, B.Q. Wei, D.H. Wu and S.W. Zhu, Mater. Sci. Eng., A271, 395–400, 1999. H. Ago, K. Petritsch, M.S.P. Shaffer, A.H. Windle and R.H. Friend, Adv. Mater., 11(15), 1281–1285, 1999. S. Kanagaraj, F.R. Varanda, T.V. Zhil’tsova, M.S.A. Oliveira, J.A.O. Sim˜oes, Composites Science and Technology, 67(15-16), 3071–3077, 2007. S. Bocchini, A. Frache, G. Camino and M. Claes, Eur Polym J, 43(8), 3222–3235, 2007. B.X Yang, J.H. Shi, K.P. Pramoda and S.H. Goh, Compos. Sci. Technol., 68(12), 2490–2497, 2008. S. Peeterbroeck, L. Breugelmans, M. Alexandre, J.B. Nagy, P. Viville, R. Lazzaroni, Ph. Dubois, Compos. Sci. Technol., 67(7–8), 1659–1665, 2007. H. Shirakawa, E.J. Louis, A.G. MacDiarmid, C.K. Chiang and A.J. Heeger, J. Chem. Soc., Chem. Commun., 16, 578–580, 1977. H. Nalwa, Handbook of Organic Conductive Molecules and Polymers, John Wiley & Sons Inc., New York, 1997. C.J. Dury, C.M.J. Mutsaers, C.M. Hart, M. Matters and D.M. Leeuw, Appl. Phys. Lett., 73, 108–110, 1998.

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131. T.F. Otero, H.S. Nalwa, Handbook of Organic Conductive Molecules and Polymers, vol. 4, John Wiley & Sons, Chichester, 1997. 132. A.F. Diaz, K.K. Kanazawa and G.P. Gardini, J. Chem. Soc., Chem. Commun., (14), 635–636, 1979. 133. K.K. Kanazawa, A.F. Diaz, W.D. Gill, P.M Grant, G.B. Street, G.P. Gardini and J.F.Kwak, Synth. Met., 1, 329–336, 1980. 134. A.O. Patil, A.J. Heeger, F. Wudl, Chem. Rev., 88, 183–200, 1988. 135. C.K. Chiang, M.A. Druy, S.C. Gau, A.J. Heeger, E.J. Louis, A.G. MacDiarmid and Y.W. Park, J. Am. Chem. Soc., 100, 1013–115, 1978. 136. J. Prokeˇs, I. Kˇrivka R. Kuˇzel, J. Stejskal and P. Kratochv´ıl, Int. J. Electronics 81, 407–417, 1996. 137. G.A. Wood and J.O. Iroh, J. Appl. Polym. Sci., 61, 519–528, 1996. 138. J.O. Iroh and G.A. Wood, J. Appl. Polym. Sci., 62, 1761–1769, 1996. 139. J.H. Park, J.M. Ko, O.O. Park and D.W. Kim, J. Power Sources, 105, 20–25, 2002. 140. S. Marchant, F.R. Jones, P.V. Wright, and T.P.C. Wong, Synth. Met., 96, 35–41, 1998. 141. E. Hakansson, A. Kaynak, T. Lin, S. Nahavandi, T. Jones and E. Hu, Synth. Met., 144, 21–28, 2004. 142. S.N. Bhadani, M. Kumari, S.K.S. Gupta and G.C. Sahu, J. App.l Polym. Sci., 64, 1073–1077, 1997. 143. R.B. Bjorklund and I. Lundstrom, J. Electron. Mater., 13, 211–230, 1984. 144. M. Miˇcuˇs´ık, M. Omastov´a, J. Prokeˇs and I. Krupa, J. Appl. Polym. Sci., 101(1), 133–142, 2006. 145. I. Sapurina, N.E. Kazantseva, N.G. Ryvkina, J. Prokeˇs, P. S´aha and J. Stejskal, J. Appl. Polym. Sci., 95, 807–814, 2005. 146. M. Miˇcuˇs´ık, M. Omastov´a, Z. N´ogellov´a, P. Fedorko, M. Trchov´a, K. Olejn´ıkov´a and I. Chod´ak, I., Eur. Polym. J., 42(10), 2379–2388, 2006. 147. P.R. Hornsby and A. Mthupha, J. Mat. Sci., 29, 5293–5301, 1994. 148. Z. Tadmor and C.G. Gogos, Principles of Polymer Processing, John Wiley & Sons Inc., NY, 1979. 149. S. Middleman, Fundamentals of Polymer Processing, McGraw-Hill, NY, 1977. 150. P.R. Hornsby, Proceedings from R&D Clearing House Meeting ‘Polymer Processing, Structure and Properties’, Leads, 15–16 May, 1991. 151. A.I. Medalia, J. Colloid. Interface. Sci., 32, 115–131, 1970. 152. W.R. Bolen and R.E. Colwell, SPE Journal, 14, 24–28, 1958. 153. Z. Tadmor, Ind. Eng. Chem. Fundam., 15, 346–348, 1976. 154. C. Klason and J. Kubat, Polymer Processing and Properties, Proceedings of the European Meeting, 47–52, 1984. 155. P.R. Hornsby, Plastics and Rubber: Processing and Applications, 7, 237–240, 1987. 156. J.W. Ess and P.R. Hornsby, Plastics and Rubber: Processing and Applications, 8, 147–156, 1987. 157. I. Chod´ak and I. Krupa, J. Mat. Sci. Lett., 18, 1457–1459, 1999. 158. J.R. Jurado, C. Moure, P. Duran, M. Rodr´ıguez, A. Linares and J.L.Acosta, J Mat Sci., 26, 4022–4025, 1991. 159. I. Chod´ak, Prog. Polym. Sci., 23, 1409–1442, 1998. 160. J. Brandrup, E.H. Immergut and E.A. Grulke (Eds), Polymer Handbook, 2 Volume set, 4th edition, Wiley-Interscience, 2003. 161. George Lubin (Ed.), Handbook of Composites, Van Nostrand Reinhold, 1982. 162. M.A. C´ardenas, D. Garc´ıa-L´opez, I. Gobernado-Mitre, J.C. Merino, J.M. Pastor, J. de D. Mart´ınez, J. Barbeta and D. Calveras, Polym. Degrad. Stab., 93(11), 2032–2037, 2008. 163. H. Liu, Q. Wu, G. Han, F. Yao, Y. Kojima and S. Suzuki, Compos. Part A: Appl. S., 39(12), 1891–1900, 2008. 164. D.R. Paul and S. Newman, Polymer Blends, vol. 1., Academic Press, NY, 1978. 165. O. Olabisi, L.M. Robeson and M.T. Shaw, Polymer–Polymer Miscibility, Academic Press, London, 1979. 166. M. Sedl´acˇ kov´a, I. Lac´ık and I. Chod´ak, Macromol. Symp., 170, 157–164, 2001. 167. L.A. Utracki, Polymer Alloys and Blends, Hanser, Toronto, 1989. 168. Yu. S. Lipatov, Ye.V. Mamunya, N.A. Gladyreva and Ye.V. Lebedev, Polym. Sci. U.S.S.R., 25(7), 1714–1722, 1983. 169. G. Geuskens, J.L. Gielens, D. Geshef and R. Deltour, Eur. Polym. J., 23(12), 993–995, 1987. 170. M. Sumita, K. Sakata, S. Asai, K. Miyasaka and H. Nakagawa, Polym. Bull., 25, 265–271, 1991. 171. M. Sumita, K. Sakata, Y. Hayakawa, S. Asai, K. Miyasaka and M. Tanemura, Colloid Polym. Sci., 270, 134–139, 1992.

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477

K. Bueppelmann and K. Weber, Plastics, Rubber and Composites, Processing and Applications, 22, 47–50, 1994. L.B. Valdes, Proc. IRE, 42, 420–427, 1954. A. Uhlir, Jr., Bell. Syst. Tech. J. 34, 105–128, 1955. M.A. Green and M.W. Gunn, Solid-St. Electron., 15, 577–585, 1972. L.J. van der Pauw, Philips Res. Rep., 13, 1–9, 1958. D.S. Perloff, Solid-State Electron., 20, 681–687, 1977. Y. Sun, J. Shi, and Q. Meng, Semic. Sci. Technol., 11, 805–811, 1996. R.G. Mazur and D.H. Dickey, J. Electrochem. Soc., 113, 255–259, 1966. T. Clarysse and W. Vandervorst, J. Vac. Sci. Technol,. B10, 432–437, 1992. T. Clarysse, I. Hoflijk, W. Zhang and W. Vandervorst, J. Vacuum Sci. Technol., B22, 444–449, 2004. R.H. Cox and H. Strack, Solid-St. Electron., 10, 1213–1218, 1967. W.D. Edwards, W.A. Hartman and A.B. Torrens, Solid-St. Electron., 15, 387–392, 1972. A. Shepela, Solid-St. Electron., 16, 477–481, 1973. C.-Y. Ting and C.Y. Chen, Solid-St. Electron., 14, 433–438, 1971. D.K. Schroder, Semiconductor Material and Device Characterization, Second Edition. John Wiley & Sons Inc., NY, 1998. 187. P. Blood and J.W. Orton, The Electrical Characterization of Semiconductors: Majority Carriers and Electron States, Academic Press, London, 1992. 188. Q. Cao and J.A. Rogers, Ultrathin Films of Single-Walled Carbon Nanotubes for Electronics and Sensors: A Review of Fundamental and Applied Aspects, Adv. Mater, 21(1), 29–53, 2009.

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12 Dielectric Spectroscopy and Thermally Stimulated Depolarization Current Analysis of Multiphase Polymer Systems Polycarpos Pissis National Technical University of Athens, Athens, Greece

Apostolos Kyritsis National Technical University of Athens, Athens, Greece

Daniel Fragiadakis Naval Research Laboratory, Washington, DC, USA

12.1 Introduction Multiphase polymer systems, such as polymer blends, copolymers, interpenetrating polymer networks (IPNs) and polymer composites, are of particular interest from both the fundamental point of view and that of applications. Different properties of the particular components, partly controversial, can be combined in such systems. Moreover, new properties, not inherent to the components, may emerge as a result of synergy [1, 2]. In addition to composition, the properties of a multiphase polymer system, including those of technological significance, depend on morphology, which can be controlled by the choice of the method and the conditions of preparation and by processing [3]. This dependence provides vast possibilities for preparing new materials with specific properties, tailored to meet specific end-use requirements. The most rational way to effectively make use of the vast possibilities offered by multiphase polymer systems is by systematically studying and better understanding the structure–property relationships. The term ‘structure–property relationship’ is used as an abbreviation of the complex relationships between composition, processing, structure/morphology, molecular dynamics and final properties. The investigation of

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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structure–property relationships is a fundamental issue in materials science. The profound understanding of these relationships provides a basis for optimizing the composition and the synthesis/processing conditions towards designing new materials with predicted properties tailored to specific applications [4, 5]. Several experimental techniques have to be combined for a systematic investigation of structure–property relationships in multiphase polymer systems. Among them, dielectric techniques are very effective and widely used, in particular in recent years, in order to investigate molecular dynamics and structure/morphology in multiphase polymer systems. In the investigation of polymer dynamics (or chain dynamics or molecular dynamics or molecular mobility, these terms are considered here as synonyms) in polymeric systems, dielectric techniques compete with, or better are complemented by, several other techniques, in particular differential scanning calorimetry (DSC) [6], dynamic mechanical analysis (DMA) [7] and nuclear magnetic resonance (NMR) [8]. Each of these techniques is characterized by special features rendering it attractive for specific applications. That is because each technique probes molecular mobility in a different way. For that reason, the combination of two or more techniques is essential, as in that case several aspects of molecular dynamics can be studied. Results for the time scale of the response by the various techniques, after appropriate analysis by methods characteristic of each technique, are often presented in common Arrhenius diagrams (activation plots) [9]. In addition to the experimental techniques mentioned above, computer simulations have attracted much interest in recent years as a promising tool for investigating chain dynamics in polymeric systems [10]. The main feature and advantage of dielectric techniques, as compared to other techniques for studying molecular dynamics in polymeric systems, is the broad frequency range, which can be relatively easily covered (10−4 –109 Hz in the examples to be presented in this chapter). A second advantage is that, in addition to polymer dynamics, charge carrier motion can be thoroughly studied by dielectric techniques. The broad frequency range of dielectric measurements, in combination with temperature variation, allows to measure on the same sample processes with very different characteristic (relaxation) times and, correspondingly, different characteristic length scales. These include fast secondary (local) relaxations at high frequencies with characteristic length scales of below 1nm; cooperative relaxations like the segmental relaxation, associated with the glass transition (dynamic glass transition [11]), and the global relaxation, related with the presence of a dipole moment component along the chain contour [12], at intermediate frequencies, with characteristic length scales of a few nm (cooperativity length of the glass transition) and of the end-to-end distance of the chain, respectively; dc conductivity and conductivity effects (like the interfacial Maxwell-Wagner-Sillars relaxation [13]) at lower frequencies, with characteristic length scales in the nm to μm range. Morphology in multiphase polymer systems, in relation to compatibility of and interactions between the components, in particular miscibility and microphase separation, is a fundamental issue. Often results reported in the literature for the same system are controversial depending on the experimental technique and the concept used. One widely-accepted criterion of miscibility in a two-component polymer system is the occurrence of a single glass transition at an intermediate temperature between the Tg ’s of the two components [14], most frequently measured by DSC, often also by DMA. However, typically broad glass transitions are measured by DSC in miscible polymer systems [14]. The broadening is, in general, discussed in terms of a distribution of Tg ’s due to composition fluctuations in the material over dimensions of the order of nanometers, giving rise to nanoheterogeneities [15]. The cooperativity length (the size of cooperatively rearranging regions in the Adam-Gibbs theory [16]) is also of the order of a few nanometers at Tg , whereas it decreases significantly with increasing temperature [11]. Thus, the use of dynamic (spectroscopic) techniques, i.e. molecular dynamics studies with variation of the frequency of the applied stimulus, to investigate the characteristics of the glass transition in miscible polymer systems offers additional possibilities for morphological studies [15, 17]. By these techniques the glass transition is probed at temperatures higher than the calorimetric Tg (dynamic glass transition [11]) depending on the frequency of measurement.

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Similar to broad glass transitions by DSC, a broadening of the spectrum of relaxation times, with respect to those of the pure components, is observed for the segmental (main) relaxation in miscible polymer systems by dynamic techniques [18]. Several models have been proposed to relate the shape of the dielectric relaxation spectrum to concentration fluctuations. It has been shown that in the case of miscible polymer blends with large difference between the Tg ’s of the two components the presence of nanoheterogeneities, resulting from concentration fluctuations, leads to the observation of two differentiated segmental relaxation processes, corresponding to the low-Tg component and to the average composition, respectively [15]. This so-called dynamic heterogeneity is caused by the decrease of the cooperativity length of the glass transition with increasing temperature [11]. Thus, at any temperature above the calorimetric Tg of the blend the length of cooperativity of the low-Tg component is very small, so that there are domains with sizes in this order of magnitude which contain only chains of this component. On the contrary, the length of cooperativity of the high-Tg component is higher than the size of the larger regions which contain only chains of this component, so that the relaxation process occurs in rearranging regions with an average blend composition [15]. Also other authors employed the concepts of concentration fluctuations and local heterogeneity, i.e. local regions rich in either one or the other component, but consider other length scales of the relevant local region [19–21]. An alternative way to interpret experimental results in miscible polymer bends with large difference between the Tg ’s of the two components was proposed by Lodge and McLeish [22, 23]. The self-concentration model proposed by these authors emphasizes the role of chain connectivity. As a direct consequence of that, the average composition of the local environment around any chosen chain segment is automatically enriched in the same species, so that each species experiences a different average local environment and each component senses its own, composition-dependent glass transition. The Kuhn length of the polymer, corresponding to the amount of backbone required for a chain to bend by 180◦ , was considered as the relevant lengthscale for evaluating self-concentration within this model [19]. The model suggests that it may be possible to understand many features of the dynamics of miscible polymer blends without invoking a cooperativity length scale that diverges at or below Tg . In particular, it provides a straightforward way to explain why a polymer blend can exhibit two glass transitions (also by DSC), even when the two components are intimately mixed [23]. Colmenero and co-workers combined an extension of the Adam-Gibbs theory [16] with the concept of self-concentration [22] to describe the temperature dependence of the relaxation time for the component segmental dynamics in miscible polymer blends. The dynamics of each component in the blend could then be obtained from the dynamics and the thermodynamics of the pure components and a single fitting parameter obtained from the fitting of the experimental data [24, 25]. It should be stressed here that for this and the other models mentioned above broadband dielectric spectroscopy was the main technique used to prove the validity of the model and/or obtain values of the fitting parameters. This chapter deals with the use of dielectric techniques to study polymer dynamics and morphology in multiphase polymer systems. In addition to the classical dielectric spectroscopy (DS) in the frequency domain, a less frequently used dielectric technique in the temperature domain, thermally stimulated depolarization current (TSDC), will be considered. This technique, known also under the names thermally stimulated depolarization (TSD) and thermally stimulated current (TSC), corresponds to measuring dielectric loss as a function of temperature at a fixed low frequency in the range 10−4 –10−2 Hz. Special attention will be paid to methodologies developed by our group in Athens and by others which make dielectric techniques a powerful tool for the investigation of polymer dynamics and morphology in multiphase polymer systems, illustrated by several examples. These methodologies are based, to a large extent, on the broad frequency range of DS, which in combination with temperature variation makes possible the investigation of processes of variable lengthscale. In the particular case of TSDC the special methodologies are based on the high sensitivity and the high peak resolving power of the technique.

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Dielectric techniques, in particular DS, have been widely used by several research groups to study various aspects of polymer dynamics and of morphology in multiphase polymer systems. Recent representative references are [26–37]. In this chapter, we focus mostly on results obtained by our research group in Athens to illustrate the concepts used and the methodologies developed and to demonstrate the power of dielectric techniques. Three systems have been selected to that end: (1) random copolymers and interpenetrating polymer networks (IPNs) of poly(alkyl acrylate)s and poly(alkyl methacrylate)s, where we focus on polymer dynamics in relation to mixing and microphase separation; (2) rubber/silica nanocomposites with a fine dispersion of silica nanoparticles, where we focus on the investigation of effects of interfacial interactions between the two components on polymer dynamics; and (3) polymer nanocomposites with conducting carbon inclusions, where we demonstrate the power of dielectric techniques to study electrical conductivity in such systems, percolation and charge transfer between conducting islands. The organization of the chapter is as follows. Dielectric techniques, in particular broadband DS and TSDC are briefly introduced in the next section. In the following sections we focus on three examples selected to demonstrate the power of dielectric techniques, copolymers and IPNs based on poly(alkyl acrylate)s and poly(alkyl methacrylate)s (Section 12.3), rubber/silica nanocomposites (Section 12.4), and polymer nanocomposites with conducting carbon inclusions (Section 12.5). The special methodologies developed for using dielectric techniques to study polymer dynamics and morphology in multiphase polymer systems will be illustrated by these three selected examples, in particular methodologies for studying mixing and phase separation at various lengthscales in Section 12.3, interfacial phenomena in Section 12.4, and percolation phenomena in Section 12.5. Final conclusions are then drawn in Section 12.6.

12.2 Dielectric Techniques 12.2.1

Introduction

The dielectric function (or dielectric permittivity or dielectric constant) ε* (ω) describes the material response to the application of an alternating electric field E(ω). For small electric field strengths a linear relationship holds between E and the polarization P: P(ω) = (ε∗ (ω) − 1)ε0 E(ω)

(12.1)

where ε0 is the permittivity in vacuum [38]. ε* (ω) is related by the theory of dielectric relaxation to the correlation function (t) of the polarization fluctuations [39–41], ε∗ (ω) − ε∞ = εs − ε∞

 ∞  d(t) exp(−iωt)dt − dt

(12.2)

0

and (t) =

P(t)P(0) P(0)2 

(12.3)

where P denotes a fluctuation of the polarization around its equilibrium value and the brackets denote the averaging over an ensemble or time t.

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The dielectric function and other dielectric properties (see below) are measured by dielectric techniques, which are a powerful tool for studying molecular dynamics in various materials, including polymers and composites. The main advantage of dielectric techniques over other techniques of measuring molecular dynamics is the extremely broad frequency range covered, which extends from about 10−5 to about 1011 Hz [39–41]. Obviously, this broad frequency range cannot be covered by a single technique and measurements are performed in the frequency domain, the time domain or the temperature domain. In most cases, measurements are performed in the frequency domain, i.e. with variation of the frequency of the applied electric field (broadband dielectric spectroscopy, DS, Section 12.2.2). In slow time domain spectroscopy (TDS), a voltage step is applied to the sample and the polarization or depolarization current is measured as a function of time. The dielectric relaxation spectrum is then obtained by Fourier transformation or approximate formulae. By carefully controlling the sample temperature and accurately measuring the depolarization current, precision measurements of the dielectric function down to 10−6 Hz are possible [42]. In fast time domain spectroscopy or reflectometry (TDR) a step-like pulse propagates through a coaxial line and is reflected from the sample section placed at the end of the line. The difference between the reflected and the incident pulses recorded in the time domain contains the information on the dielectric properties of the sample. By transformation in the frequency domain the dielectric function is obtained, typically in the frequency range of about 10 MHz–10 GHz [43, 44]. Finally, a special dielectric technique in the temperature domain is that of thermally stimulated depolarization currents (TSDC, Section 12.2.3). 12.2.2

Broadband Dielectric Spectroscopy (DS)

In most cases measurements are carried out isothermally in the frequency domain and the term (broadband) dielectric spectroscopy (DS) or (broadband) dielectric relaxation spectroscopy (DRS) is then used. Other names frequently used for DRS are impedance spectroscopy and admittance spectroscopy. Impedance spectroscopy is usually used in connection with electrical conductivity, electrolytes and electrochemical studies, whereas admittance spectroscopy often refers to semiconductors and devices. For measurements in the frequency domain, capacitance bridges, impedance analysers, frequency response analysers, radio-frequency reflectometers and network analysers are typically employed. In Figure 12.1 we

10–6 10–4 10–2

100

102

104

106

108

1010 1012

f (Hz)

FRA AC bridges Imped. anal. Q-meters network analyzers slow TDS TDS

frequency domain

time domain

Reflectometers Resonance circuits Cavities and waveguides

Figure 12.1 Techniques and equipment for dielectric measurements. FRA means frequency response analyzer; TDS is time domain spectroscopy.

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show schematically the frequency range of dielectric measurements covered by different techniques and equipment [41]. The principle of these measurements is as follows. The sample under investigation is placed in a capacitor with empty capacitance C0 , which becomes a part of an electric circuit. A sinusoidal voltage with angular frequency ω is applied to the circuit and the complex impedance Z * (ω) of the sample is measured. The complex dielectric function ε* (ω) = ε (ω)-iε (ω), defined by ε∗ (ω) =

C∗ C0

(12.4)

where C is the capacitance of the filled capacitor, is then obtained from ε∗ (ω) =

1 iωZ ∗ (ω)C

(12.5) 0

Independently of the specific dielectric technique used, the results of dielectric measurements are usually analyzed in the form of complex dielectric function ε∗ (ω) = ε (ω) − iε (ω) at constant temperature by fitting empirical relaxation functions to ε* (ω). In the examples to be given later in this chapter, often the two-shape-parameters Havriliak-Negami (HN) expression [45] ε∗ (ω) − ε∞ =

ε [1 + (iωτ )1−α ]β

(12.6)

is fitted to the experimental data for a relaxation mechanism. In this equation ε is the dielectric strength, ε = εs − ε∞ , where εs and ε∞ are the low- and high-frequency limits of ε , respectively, τ is the relaxation time, τ = 1/2π f HN , where f H N is a characteristic frequency closely related to the loss peak frequency f max , and α, β are the shape parameters describing the shape of the ε (ω) curve below and above the frequency of the peak, respectively, 0 < (1 − α) ≤ 1 and 0 < (1 − α)β ≤ 1. This expression becomes the single Debye form for α = 0, β = 1, the symmetric Cole-Cole form for α = 0, β = 1, and the asymmetric Cole-Davidson form for α = 0, β = 1 [45]. A proper sum of HN expressions is fitted to ε* (ω) in the case of more than one overlapping mechanisms plus a term for the contribution of conductivity, if the latter makes a contribution at the temperature of measurements. For each relaxation mechanism there are then three sources of information: the timescale of the response (τ or f max ), the dielectric strength (ε) and the shape of the response (α, β). By measuring ε* (ω) at several temperatures, the time scale of the response is analyzed in terms of the Arrhenius equation for secondary relaxations and the Vogel-Tammann-Fulcher (VTF) equation for the primary α relaxation and valuable information on the activation parameters is obtained [40, 41]. Examples will be given later in this chapter. In addition to the dielectric function ε* , which is the most physically meaningful dielectric variable to describe the material response, under the condition that in the experiments the electric field is the independent variable and the charge is the dependent one (i.e. ε is a compliance), other quantities and corresponding formalisms are often employed to analyze relaxation phenomena, in particular in systems where conductivity makes a significant contribution: electric modulus, impedance, conductivity. The (complex) electric modulus is defined by [46]: M∗ =

1 ε ε   = M + i M = + i ε∗ ε2 + ε2 ε2 + ε2

(12.7)

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The formalisms of ε* (ω) and M * (ω) are equivalent. Transformation from one to another may emphasize and, therefore, help resolve particular aspects of the relaxation process, but no new information can be extracted. In particular, space charge effects are suppressed in the modulus formalism and effects related with conductivity can be more clearly studied. In the complex impedance formalism Eq. (12.5) is used for transformation. Plots of Z * (ω) against ω provide then information on the electrical properties of the material under investigation [47]. Very popular and comprehensive are plots of -Im(Z * ) against Re(Z * ) with the frequency ω as parameter (complex impedance plots), which provide (often by extrapolation) directly the value of resistance and, by knowing the geometrical characteristics of the sample, the value of dc conductivity σ dc at the temperature of measurements. Complex impedance plots are also suitable for separation of bulk from surface (interfacial) phenomena [47]. When the emphasis is put on the bulk electrical properties of a conducting material, in comparison with the predictions of theoretical models, the ac conductivity formalism may be used [48, 49]. In that case ac conductivity plots, i.e. plot of ac conductivity σ ac against frequency ω at constant temperature, are well suited for presenting and discussing the results of conductivity measurements. The data are recorded isothermally ∗  (ω) is calculated from these data (in fact real part σac of the complex with variation of the frequency ω and σac conductivity), e.g. in the admittance presentation:  (ω) = ωε0 ε (ω) σac

(12.8)

Jonscher suggested the following power law dependence (universal dynamic or dielectric response) [48]:  (ω) = σdc + Aωs , 0.5 < s < 0.7 σac

(12.9)

where A and s are temperature dependent parameters. Equation (12.9) and modifications of that equation [49] are frequently used for fitting to the experimental data. Examples will be given later in this chapter. 12.2.3

Thermally Stimulated Depolarization Current (TSDC) Techniques

The thermally stimulated depolarization current (TSDC) method, known also under the names thermally stimulated depolarization (TSD) and thermally stimulated current (TSC), is a dielectric method in the temperature domain, which allows for a fast characterization of the dielectric response of the material under investigation. The method consists of measuring the thermally activated release of stored dielectric polarization. It corresponds to measuring dielectric losses against temperature at constant low frequencies of 10−2 –10−4 Hz [50, 51]. The low equivalent frequency is a characteristic feature of the TSDC method, which is often used to extend the range of dielectric measurements down to low frequencies. By the TSDC method, the sample is inserted between the plates of a capacitor and polarized by the application of an electric field Ep at temperature Tp for time tp , which is large in comparison with the relaxation time at Tp of the dielectric dispersion under investigation. With the electric field still applied, the sample is cooled to temperature T 0 (which is sufficiently low to prevent depolarization by thermal excitation) and then is short-circuited and reheated at a constant rate b. A discharge current is generated as a function of temperature, which is measured with a sensitive electrometer. The resultant TSDC spectrum typically consists of several peaks whose shape, magnitude and location provide information on the time scale and the dielectric strength of the various relaxation mechanisms present in the sample [50]. In the case of a single relaxation process obeying the Arrhenius equation: τ (T ) = τ0 exp

w kT

(12.10)

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the depolarization current density J(T) is given by the equation: ⎛ ⎞ T  w  w 1 P0 exp ⎝− dT  ⎠ J (T ) = exp − exp − τ0 kT bτ0 kT

(12.11)

T0

where τ is the relaxation time, w the activation energy of the relaxation, τ 0 the pre-exponential factor, T the temperature, k Boltzmann’s constant and P0 the initial polarization. The activation energy w is calculated from the measured TSC thermogram typically by the initial rise method, by fitting Eq. (12.11) to the thermogram and by simple, approximate formulae [50, 52]. The initial rise method is based on the fact that the integral term in Eq. (12.11) is negligible at T much smaller TM , where TM is the temperature of the depolarization current maximum, so that ln J (T ) ≈ cons tan t −

w kT

(12.12)

Having calculated w, τ 0 is obtained from the equation τ0 =

 w kTM2 exp − bw kT

(12.13)

The contribution of a TSDC peak to the static dielectric permittivity (relaxation strength) ε is obtained by: ε =

Q Aε0 E p

(12.14)

where A is the cross-sectional area of the sample and Q the depolarization charge evaluated from the area enclosed under the peak [50]. The TSDC method is characterized by high sensitivity and, owing to its low equivalent frequency of 10−2 –10−4 Hz, by high resolving power. In addition, it provides special variants to experimentally analyze complex relaxation mechanisms into approximately single responses: partial polarization, partial heating analysis, thermal sampling [50, 51, 53]. Examples will be given later in this chapter. An additional advantage of TSDC is based on the fact that, in contrast to DS (isothermal measurements in the frequency domain), the stages of polarization and depolarization (stimulus and response) are separated in the TSDC method. This is beneficial with respect to conductivity effects in electrically conducting polymers, where dipolar processes (typically the segmental relaxation associated with the glass transition) are often masked by ionic conductivity in DS measurements; however, not in TSDC measurements [53].

12.3

12.3.1

Copolymers and Interpenetrating Polymer Networks Based on Poly(alkyl acrylate)s and Poly(alkyl methacrylate)s (Mixing and Phase Separation) Introduction

Polyacrylate-polymethacrylate pairs are immiscible except for very low molecular weight polymers. However, by using the concept of forced compatibilization, single phase systems can be prepared in the form of sequential interpenetrating polymer networks (IPNs) and random copolymers [54]. An IPN may be defined as

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a combination of two or more polymers in network form, at least one of which is synthesized or crosslinked in the immediate presence of the other. A sequential IPN is formed by swelling the polymer network of the first component in the monomer of the second component and then polymerizing the second component. Forced compatibilization of the immiscible IPN components may be achieved in sequential IPNs if the crosslinking density of the component polymerized first is high enough, so that the second network is forced to grow in confined space by interpenetrating the first network [54]. In random copolymers, on the other hand, forced compatibilization is achieved by the covalent bonding of the monomers of the two components. In a series of papers sequential IPNs with low and high crosslinking density and random copolymers of various poly(alkyl acrylate)s and poly(alkyl methacrylate)s were prepared and investigated with respect to mixing and phase separation [54–59]. These materials combine properties of the two components and are interesting for various applications, including scaffolds for tissue engineering when one of the two components is hydrophilic, like poly(hydroxyethyl acrylate) (PHEA) [53, 58, 59]. In that case the second hydrophobic component provides the required mechanical stability. It is interesting to note in this connection that, in order to further improve mechanical stability by keeping all the other good properties, in particular hydrophilicity and biocompatibility, recently nanocomposites were prepared and investigated on the basis of these IPNs or copolymers and silica generated by sol-gel processes simultaneously with polymerization [60]. Dielectric techniques, broadband DS and TSDC, and dynamic mechanical analysis (DMA) were the main techniques employed to study mixing and phase separation at various lengthscales in the poly(alkyl acrylate)spoly(alkyl methacrylate)s IPNs and copolymers. These polymers exhibit the main α relaxation, which arises from the cooperative conformational rearrangements of the segments of the polymer chains and is associated with the glass transition, and the secondary β relaxation, arising from the rotation of the side chain with respect to its bond with the main chain. Both relaxations are mechanically and dielectrically active, the dielectric activity coming from the permanent dipole moment of the carbonyl group of the side chain. The two relaxations are interdependent and merge at temperatures higher than Tg in a new αβ relaxation process, as schematically indicated in Figure 12.2. The crossover region is the frequency–temperature region in which the relaxation times of the α and β processes approach each other (regime b in Figure 12.2). There is a rich literature and an ongoing discussion on the interdependence of the two relaxations and their merging, which might be a general phenomenon in amorphous polymers [61–63]. A key question in this discussion refers to the degree of cooperativity of the molecular motions that are responsible for the different relaxation processes. One way to follow this question is by studying the influence of a second component on the mobility of the polymer chains, in the form of IPNs of low and high crosslinking density and of copolymers in this section.

log f

a

αβ

b

c

β β α

1/T

Figure 12.2 Scheme of an Arrhenius diagram showing the merging of the α and β relaxations of a poly(alkyl methacrylate). Reprinted from [56]. Copyright (2004) with permission from American Chemical Society.

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Poly(butyl acrylate)-Poly(butyl methacrylate) Sequential Interpenetrating Polymer Networks

Two series of sequential IPNs of poly(butyl acrylate) (PBA) and poly(butyl methacrylate) (PBMA), PBA-iPBMA, with 0.1 and 10 wt%, respectively, ethylene glycol dimethacrylate (EGDMA) as crosslinking agent were prepared by block polymerization. The PBA network was polymerized first. Details of preparation have been given in [54]. The IPN samples are coded by IPNXX-YY, XX indicating the weight percent of PBA network in the IPN and YY that of EGDMA. DMA measurements show, apart from details, a single main relaxation at a temperature between those of the pure component networks in the highly crosslinked IPNs, whereas in the IPNs with 0.1% EGDMA the main relaxations of the PBA and PBMA networks are clearly separated from each other. These results suggest phase homogeneity in the first case and phase separation in the second [54]. In what follows we focus on the results and the analysis by broadband DS and TSDC. As an example of isothermally recorded DS data Figure 12.3 shows the frequency dependence of dielectric loss for IPN36-10 at several temperatures [55]. Similar to DMA, only one broad loss peak is observed which, with increasing temperature, shifts to higher frequencies, increases in magnitude and becomes narrower. The corresponding plot of the real part of the dielectric function ε (f), not shown here, exhibits a step in the region of the loss peak. In the IPNs with low crosslinking density, on the other hand, two relaxation regions are observed, corresponding to the two network components. Similar measurements to those shown in Figure 12.3 were performed with the different compositions and the pure network components to follow the β relaxation at low temperatures, the splitting into α and β at intermediate temperatures and the αβ relaxation at higher temperatures [54]. Differences are observed between the behavior by DS and by DMA due to the different relative intensity of the β and α relaxation (β relatively stronger in DS) and, more significantly, the broader frequency range of DS measurements, so that in dielectric measurements the main relaxation is measured at temperatures well apart from the calorimetric Tg (dynamic glass transition). Before we proceed with analysis of the dielectric data by fitting appropriate model functions (Section 12.2.2), we would like to extract the maximum amount of information from the raw data. In Figure 12.4 we

Figure 12.3 Dielectric loss ε against frequency f in the PBA-i-PBMA IPN36-10 at several temperatures between 30 and 120◦ C in steps of 10◦ C. Reprinted from [5]. Copyright (2001) with permission from John Wiley & Sons.

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Figure 12.4 Temperature dependence of the complex dielectric function measured at 1 kHz in PBA-i-PBMA materials: (a) ε in PBA10 network (), IPN81-10 (), IPN67-10 (◦), IPN56-10 (䉬), IPN45-10 (), IPN32-10 () and PBMA10 network (•) and (b) ε  in PBA10 network (), IPN81-10 (•), IPN67-10 (), IPN56-10 (), IPN45-10 (♦), IPN32-10 (+) and PBMA network (). Reprinted from [54]. Copyright (2001) with permission from American Chemical Society.

show isochronal (constant frequency) plots of the real and the imaginary part of the dielectric function against frequency for the highly crosslinked IPNs at 1 kHz. The data have been recorded isothermally (Figure 12.3) and replotted here. There are certain similarities to the corresponding DMA and TSDC plots (not shown here), the main difference being that in DS the frequency of presentation can be varied in a wide range of several decades. Plots at different frequencies utilize specific advantages of the technique, such as higher resolving power at lower frequencies and suppression of conductivity at higher frequencies. The ε curves in Figure 12.4b show the main α dielectric relaxation of PBA and PBMA merging into a single peak in the IPNs, clearly revealing that it consists of two overlapped relaxation processes. Analysis will enable later to separate these processes. In the corresponding plot for the loosely crosslinked IPNs, not shown here, two separated α peaks are observed whose temperatures are slightly dependent on the composition of the IPN. It should be stressed, from the methodological point of view, that the plots in Figure 12.4a enable to directly determine the relaxation strength of the corresponding processes. We will come back to this point later. The dielectric data, similar to those shown in Figure 12.3 for IPN36-10, were analyzed by fitting the HN model function (Eq. (12.6)) or a sum of two HN terms plus a conductivity term (see below), depending on the sample and the temperature of measurements [54, 55]. Analysis provides the number of relaxations contributing to a peak (one or two in this case) and for each relaxation information on the time scale, the strength and the shape. In the following we focus on the highly crosslinked networks and discuss time scale, strength and shape of the relaxations. Time scale is best discussed on the basis of the Arrhenius plot (activation diagram) shown in Figure 12.5 for all the compositions. For most of them the loss peaks have been considered as single peaks, f max in Figure 12.5 being the frequency of maximum loss. For the compositions IPN56-10 and IPN67-10 two loss peaks have been observed in the temperature and frequency range of measurements, so at several temperatures a sum of two HN terms has been fitted to the data and two values of f max are given in Figure 12.5. The data for pure PBA-10 and the compositions IPN81-10 and IPN67-10 (high frequency peak for the latter) show the typical

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Figure 12.5 Arrhenius plot of the DS data in PBA10 (), IPN81-10 (*), IPN67-10 (䉬), IPN56-10 (), IPN45-10 (+), IPN36-10 () and PBMA10 (raw data •, and data obtained by analysis ◦, details in text). The lines are fittings of the VTF equation (12.15) to the PBA-10 data and of the Arrhenius equation (12.16) to the IPN36-10 data. Note that two relaxations are observed for the compositions IPN56-10 and IPN67-10, see the text for details. Reprinted from [5]. Copyright (2001) with permission from John Wiley & Sons.

behavior of dynamic glass transition. The Vogel-Tammann-Fulcher (VTF) equation [11, 54] :  f max = A exp −

B T − T0

(12.15)

where A, B and T 0 (Vogel temperature) are temperature-independent empirical parameters, was fitted to the data for these three samples and reasonable values of the fitting parameters were obtained. An example of fitting is shown in Figure 12.5 for the pure PBA network, for which the fitting parameters take the values A = 1.4 × 1013 Hz, B = 1969 K and T 0 = 174.5 K. For the pure PBMA network the raw dielectric data (ε (f ) at several temperatures, not shown here) suggest, in agreement with data in the literature [61, 62], a splitting of the single αβ relaxation at high temperatures (or lower frequencies) into two relaxations, α and β, at lower temperatures (Figure 12.2). The results of analysis are also shown in Figure 12.5. At temperatures higher than about 120◦ C the experimental ε (f ) data can be reproduced with a single HN process, the αβ process (regime a in Figure 12.2). The intensity of the relaxation decreases with increasing temperature in this region (Figure 12.6). For T ≤ 110◦ C a sum of two HN functions gives better fits than one, α and β relaxations (regime b in Figure 12.2). In this regime the relaxation strength of the β process increases with temperature, whereas that of the α process decreases tending to zero at around 120◦ C (onset temperature of the α relaxation [54, 61]). With the temperature decreasing from 110 to 70◦ C, the shape parameters, provided by the HN analysis, decrease. For the α peak (1 − α) decreases from 0.64 to 0.56 and (1 − α)γ from 0.35 to 0.23, whereas for the β peak (1 − α) decreases from 0.48 to 0.42 and (1 − α)γ from 0.34 to 0.26. We recall that the shape parameters (1 − α) and (1 − α)γ give the slope of ε (f ) at frequencies lower and higher, respectively, than the peak frequency (Section 12.2.2). It follows that both

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Figure 12.6 Relaxation strength of the components α (•) and β () of the relaxation spectrum of PBMA10 obtained by fitting the experimental curves to the sum of two Havriliak–Negami distributions. Two points () corresponding to the αβ relaxations are also included (see text).

peaks are asymmetrical (steeper at lower frequencies) and become broader with decreasing temperature and that at each temperature the β peak is broader than the α peak [55, 62]. It is interesting to note from the methodological point of view that information on the relaxation strength is provided also by the raw data without recourse to curve fitting, namely the comparative ε (T) plot for several compositions in Figure 12.4 and the Cole-Cole diagram for PBMA-10, as an example, in Figure 12.7. The latter is a plot of ε against ε with the frequency f as parameter and summarizes data taken isothermally at several temperatures. We observe in Figure 12.7 that the limit of the isotherms at low ε values (which corresponds to high frequencies) is nearly temperature independent, whereas the high value limit (which corresponds to low frequencies) clearly decreases with increasing temperature above 100◦ C. The dielectric data in Figure 12.5 for the compositions IPN36-10, IPN45-10, IPN56-10 and IPN 67-10 (the low-frequency relaxation for the last two compositions) suggest a linear dependence with a slope similar to that for the β relaxation in the pure PBMA network. As an example, a fit of the Arrhenius equation [11, 55]  w f max = f 0 exp − kT

(12.16)

where f 0 a constant, w the activation energy and k Boltzmann’s constant, to the IPN36-10 data has been included in Figure 12.5. The fitting parameters are w = 0.90 eV and f 0 = 7.4 × 1016 Hz. These results suggest a suppression of cooperativity (giving rise to the VTF behavior) in the IPNs and may be discussed in terms of constraints imposed by crosslinking, entanglements, and confinement [64]. On the other hand, with increasing PBA content, f max shifts, at constant frequency, to lower temperatures towards PBA, like the α relaxation should do. In addition, the values of w and f 0 are rather high for one-barrier activation, as commonly assumed

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Figure 12.7 Cole–Cole diagrams of PBMA10 network: () 30◦ C, (•) 40◦ C, () 50◦ C, () 60◦ C, (♦) 70◦ C, (+) 80◦ C, () 90◦ C, (◦) 100◦ C, () 110◦ C, () 120◦ C, (♦) 130◦ C, (×) 140◦ C, (*) 150◦ C. Reprinted from [54]. Copyright (2001) with permission from American Chemical Society.

for local relaxations. These results suggest a change with composition, from the splitting of the single αβ relaxation into two relaxations, α and β, in PBMA (Figure 12.2) to a Goldstein-Johari β relaxation [65] in the densely crosslinked IPNs. Similar results were obtained with highly crosslinked sequential IPNs of poly(ethyl acrylate) (PEA) and poly(ethyl methacrylate) (PEMA): splitting of the single αβ relaxation into two relaxations, α and β, in PEMA, a single α relaxation in PEA-rich compositions and Arrhenius behavior in middle compositions [57]. These results were explained by suppression of cooperativity due to incorporation of PEA chains to the cooperative motions of the PEMA network, which shifts the α and β relaxations to a different extent, inducing changes in the shape of the spectrum in the merging zone [57]. It is interesting to note that in nanoheterogeneous polycyanurate networks modified by poly(oxytetramethylene glycol) a change of dynamics from a cooperative α relaxation to a Goldstein-Johari beta relaxation was observed with decreasing temperature [66] and discussed in terms of the Adam-Gibbs theory for the glass transition [16]. The suppression of cooperativity in these systems was quantified in terms of fragility. The concept of fragility, i.e. deviation from the Arrhenius behavior in the activation diagram (Figure 12.5), introduced by Angell [67], has been much used in recent years to classify glass-forming materials with respect to kinetic and thermodynamic aspects of the glass transition. Fragility in these systems, calculated from the dielectric data [66, 67], decreases with inceasing content of poly(oxytetramethylene glycol), i.e. cooperativity is suppressed. A similar quantification of the suppression of cooperativity in the IPNs studied here is less straightforward because of the complex splitting scenario of Figure 12.2 in PBMA. Coming back to the activation diagram in Figure 12.5, two relaxations are observed by DS for the compositions IPN56-10 and IPN67-10, one of them with characteristics similar to those of the α relaxation in PBA and the other corresponding to the mean composition. These results may be discussed in terms of dynamic heterogeneity and the model proposed by Kumar et al. [15, 55]. Figure 12.8 compares characteristic temperatures of the dynamic processes in the highly crosslinked IPNs determined by mechanical and dielectric techniques: the temperature of the DMA tanδ peak at 1 Hz and the temperature of the DS loss peak at 1 kHz (Figure 12.4) (the low frequency peak for the compositions IPN56-10 and IPN67-10). For the PBA rich samples two more temperatures have been included: that of the

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Figure 12.8 Temperature of DMA tanδ peak at 1 Hz (), temperature of DS loss peak at 1 kHz (◦), temperature of TSDC peak () and Tg diel ( ) vs composition of the highly cross-linked PBA-i-PBMA IPNs. The line is the Fox equation (12.17), details in text. Reprinted from [5]. Copyright (2001) with permission from John Wiley & Sons.

TSDC α peak and T g diel , defined by the condition f max (T g diel ) = 1.6 × 10−3 Hz, corresponding to a relaxation time of 100 s [12] and obtained by extrapolation of the data in Figure 12.5 to low frequencies. From the methodological and the metrological point of view it should be stressed that the two measures of the glass transition temperature Tg provided by dielectric techniques, the temperature of the TSDC α peak and T g diel , are typically very close to the calorimetric Tg [9, 12]. All the characteristic temperatures show the same trend of increasing with decreasing PBA content. For a detailed discussion the time scale, defined by the frequency of measurement, and the spatial scale, given by the type of motion probed, should be considered. In addition, it should be taken into account that a modulus is measured by DMA, whereas a compliance by DRS and that the characteristic temperatures refer to tanδ for DMA, whereas to ε for DS [9]. The line in Figure 12.8 is the Fox equation [55] for a homogeneous mixture, 1 wPBA (1 − wPBA ) = + T TPBA TPBMA

(12.17)

where TPBA and TPBMA the characteristic DMA temperatures of the pure networks, which coincide with the DS (1 kHz) temperatures. It describes well the DMA and the DS (1 kHz) data, providing additional support for miscibility achieved by forced compatibilization [55]. 12.3.3

Poly(butyl acrylate)-Poly(methyl methacrylate) Interpenetrating Polymer Networks and Copolymers

DMA, broadband DS and TSDC were used in a complementary way to study the merging of the main α and the secondary β relaxations in poly(butyl acrylate)-i-poly(methyl methacrylate) (PBA-i-PMMA) sequential IPNs in comparison to net-poly(butyl acrylate)-co-poly(methyl methacrylate) (net-PBA-co-PMMA) random copolymer networks [56]. Both IPNs and copolymer networks were highly crosslinked, so that forced

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compatibilization was achieved also in the IPNs (compare Section 12.3.2), as shown by DMA and TSDC. Broadband DS provided a detailed picture of polymer dynamics by making use of the fact that in the dielectric spectra the β relaxation dominates over the α relaxation in PMMA, whereas the opposite is the case in PBA. In the copolymer networks the introduction of BA segments in the polymer chains shifts the α process towards lower temperatures with respect to PMMA and increases the intensity of the α relaxation. The crossover is shown in copolymers with MMA content between 20 and 60 wt%; however, not in the IPNs with similar compositions, where the spectra show the characteristics of a secondary relaxation. This result is explained by the spreading of the glass transition in an extremely broad temperature range in the nanoheterogeneous IPNs, so that at each temperature the intensity of the α relaxation is small compared to that of the β relaxation. The study demonstrated the power of dielectric techniques to reveal fine differences in polymer dynamics in the two multiphase systems both macroscopically homogeneous and with the same composition. Here we focus only on two aspects of the use of dielectric techniques not illustrated in the previous section, namely reduced Cole-Cole plots in relation to dielectric strength and shape of the response and TSDC thermograms for a quick characterization of the overall dielectric behavior. Preparation of the IPNs and the copolymers has been described in [56]. IPNs are designated as IPNXX, XX being the weight percentage of net-PBA in the IPN. Copolymer networks are designated as COPZZ, ZZ being the weight percentage of BA units in the copolymer. Raw dielectric ε (f ) data taken at several temperatures indicate, and analysis by the HN model function confirms, that in the PBA network the response is dominated by the alpha relaxation and the shape of the ε (f ) isotherms does not change with temperature, indicative of a thermorheologically simple behavior. Reduced Cole-Cole diagrams, where ε /ε is plotted against (ε −ε∞ )/ε, help to further clarify this point. Figure 12.9a shows that plot for the PBA network [56]. The Cole-Cole arcs for the different temperatures superimpose on a master curve. Thus, the shape of the relaxation is temperature-independent, in agreement with the results of the HN analysis giving values of α = 0.63 and γ = 0.54 independent of temperature. Please note, however, that reduced Cole-Cole diagrams are solely based on raw data without any analysis, similar to the Cole-Cole diagrams of Figure 12.7. Also for COP80 the normalized Cole-Cole arc is a master curve and the values of the shape parameters α = 0.7 and γ = 0.55 are

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Figure 12.9 Reduced Cole–Cole diagrams for the polymer and the net-PBA-co-PMMA copolymer networks. (a) net-PBA: () −30, () −25, (•) −20, () −15, (◦) −10, ( ) −5 and () 0◦ C. (b) COP60: ( ) 10, (◦) 20, () 30, () 40, () 50, and (•) 60◦ C. Reprinted from [56]. Copyright (2004) with permission from American Chemical Society.

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0 50 100 Temperature (°C)

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Figure 12.10 TSDC thermograms obtained with the homopolymer and the net-PBA-co-PMMA copolymer networks and two PBA-i-PMMA IPNs indicated on the plot: () net-PMMA, () COP20, (䉬) COP40, () COP60, (•) COP80, (◦) net-PBA, ( ) IPN42, and () IPN59. Only one datum per degree has been represented for clarity. Reprinted from [56]. Copyright (2004) with permission from American Chemical Society.

temperature-independent. No master curve could be obtained for the rest of the copolymers, as shown in Figure 12.9b for COP60, and for the PMMA network, these results reflecting the increasing influence of the secondary β relaxation of PMMA with shape and dielectric strength depending on temperature. In agreement with that the shape parameter α decreases with temperature for these samples, whereas γ is approximately constant. Figure 12.10 shows comparative TSDC thermograms for the homopolymer and the copolymer networks and two IPNs in the temperature region of interest for each network [56]. From the methodological point of view we would like to stress the simplicity and power of the method which provides a quick characterization of the overall dielectric behavior of a material, as illustrated in Figure 12.10. The TSDC thermograms can be compared with DMA thermograms, recorded also with increasing temperature, and isochronal ε (T) plots at selected fixed frequencies [56]. Two relaxations are observed in Figure 12.10 in net-PMMA, the α relaxation at about 140◦ C and the broad β relaxation at about –40◦ C, whereas net-PBA exhibits only one relaxation, the α relaxation at about –40◦ C. The copolymers show also two relaxations: a broad α relaxation systematically shifting to higher temperatures with decreasing PBA content and a weaker relaxation in the temperature region of the β relaxation of net-PMMA, which for COP80 and COP60 becomes a shoulder on the low-temperature side of the α relaxation. The TSDC method provides special techniques for experimentally analyzing such complex relaxations into approximately single responses. We will provide examples for that in Section 12.4. Similar to the copolymers, two peaks are observed in the thermograms of the IPNs. The high-temperature peak is shifted toward the PMMA α peak, with respect to the corresponding copolymer peak, whereas the low-temperature peak looks more similar to the PBA α peak rather than the PMMA β peak, both effects being more clear for IPN59 than for IPN42. We refer to [56] for further details on the significance of these results for the materials and the topics under investigation and their correlations to results obtained by DS and by DMA.

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Poly(ethyl methacrylate)-Poly(hydroxyethyl acrylate) Copolymers

The segmental dynamics of poly(ethyl methacrylate-co-hydroxyethyl acrylate), P(EMA-co-HEA), random copolymers was analyzed in [58] using DMA, TSDC and broadband DS. These materials are interesting with respect to applications as hydrogels with improved mechanical stability provided by the hydrophobic PEMA (Section 12.3.1). The results showed that the distribution of relaxation times, analyzed by the HN model function (6), is broader in the copolymers with intermediate composition than in the homopolymers, a feature that can be explained by the inhomogeneity produced on the molecular scale by the sequence distribution of the monomeric units along the chain. In consistency with that the copolymers exhibit a lower kinetic fragility than the simple average of those of the homopolymers [58]. Here we illustrate at this example the power of broadband DS and TSDC by focusing on aspects of the methods not particularly stressed in Sections 12.3.2 and 12.3.3. Details of the preparation of the copolymers, designated as P(EMA-co-HEA)XX/YY, XX and YY being the weight fraction of EMA and HEA units, respectively, are given in [58]. Figure 12.11 shows normalized TSDC thermograms for all the compositions studied. Normalization refers to the intensity calculated by In = Id/Vb, where I is the measured depolarization current, b the heating rate, V the polarizing voltage and d the thickness of the sample. Normalized TSDC thermograms allow for a direct comparison of relaxation strengths of the relaxations recorded by using Eq. (12.14) [50, 51]; however, conductivity and space charge contributions have to be properly taken into account (see below). The thermogram for pure PEMA, the most hydrophobic polymer of the series, presents, in the order of increasing temperature, the secondary β relaxation, the main α relaxation around 75◦ C and the ρ peak related with space charge relaxation around 100◦ C. As the content of HEA units in the copolymers increases, the contribution of conductivity and of the space charge relaxation increases, the ρ peak can no longer be observed as a separate peak and the dielectric strength of the α peak, which shifts systematically to lower temperatues, can no longer be reliably determined. Nevertheless, the temperature of the α peak, which is a good measure of Tg , can be accurately determined, as demonstrated in a detailed study on PHEA hydrogels [53]. In the same study a methodology was illustrated for separating the dipolar α peak from conductivity and space charge contributions, based on the use of highly insulating foils between the sample and the electrodes [53]. A final comment with respect to Figure 12.11 refers to the fact that the TSDC α relaxation can be reliably recorded even in pure PHEA, whereas that is not possible by DS, due to increased conductivity [68]. We commented on that possibility in Section 12.2.3. As mentioned above, TSDC provides easily accurate values of the temperature of the α peak, T α . Figure 12.12 shows T α as a function of composition, together with two more measures of Tg , the temperature of the ε (T) loss peak determined by DS at a high frequency of 100 kHz (to eliminate/suppress effects of conductivity) and the temperature where the mechanical relaxation time determined from DMA master curves becomes 100 s (corresponding to a frequency of 1.6 mHz) [58]. All three measures of Tg show the same trend with composition, well described by the Fox equation (12.17) for miscible blends. The different absolute values reflect the influence of the frequency of measurements, e.g. DS-TSDC where the motions probed should be similar but the corresponding frequencies very different (100 kHz in DS against the equivalent frequency of TSDC in the range of 10−2 –10−4 Hz [50, 51]), and of the type of motions probed, e.g. TSDCDMA where frequencies should be rather comparable but the type of motions probed may differ significantly [9, 39–41]. Figure 12.13 shows dc conductivity as a function of temperature for the homopolymers and four copolymers. Equation (12.8) was used to calculate ac conductivity from the measured ε (f ) with temperature as parameter and then dc conductivity was obtained as the plateau value from ac conductivity plots σ ac (f ). We will make extensive use of the ac conductivity formalism later in Section 12.5. The investigation of dc conductivity may be utilized for morphological characterization, as the moving ions probe the local morphology [58]. Two effects contribute to the increase of the conductivity at each temperature with increasing PHEA content

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Figure 12.11 Normalized TSDC thermograms measured for all the samples indicated on the plot. The polarization temperatures are: 90◦ C for PEMA, P(EMA-co-HEA)80/20 and P(EMA-co-HEA)60/40, 40◦ C for P(EMAco-HEA)50/50 and 30◦ C for P(EMA-co-HEA)80/20 and PHEA. Reprinted from [58]. Copyright (2007) with permission from Elsevier.

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Figure 12.12 Temperatures of main relaxation of the P(EMA-co-HEA) copolymer as a function of the weight fraction of HEA units in the copolymer network. ( ) Maximum of the TSDC peak corresonding to the α relaxation, (◦) temperature of the ε peak at 105 Hz, (•) glass transition temperature determined from the dynamic-mechanical muster curves as the temperature at which the mechanical relaxation time equals 100s. The solid lines represent the Fox result. Reprinted from [58]. Copyright (2007) with permission from Elsevier.

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Figure 12.13 dc conductivity for homopolymer and copolymer networks: (*)P(EMA-co-HEA)60/40, () P(EMA-co-HEA)50/50, (•) P(EMA-co-HEA)40/60, (◦) P(EMA-co-HEA)20/80 and () PHEA. Reprinted from [58]. Copyright (2007) with permission from Elsevier.

shown in Figure 12.13 [68]. The first is a temperature effect, related to the facts that (a) at temperatures above the glass transition temperature Tg, in the liquid state, conductivity increases with the difference T-Tg , and (b) Tg decreases with the PHEA content of the copolymer. However, the temperature effect does not explain completely the increase in conductivity shown in Figure 12.13, as can be proved representing σ dc against T-Tg or against T /Tg (Tg -normalized Arrhenius plot [68], results not shown here). The other effect is the higher conductivity of PHEA as compared to PEMA owing to a higher concentration of charge carriers, for which traces of water may play an important role [68]. The results in Figure 12.13 show that the samples can be classified into two groups. In the first group (P(EMA-co-HEA)60/40 and P(EMA-co-HEA)50/50) the samples are characterized by low values of dc conductivity and weak temperature dependence. In the samples of the second group (P(EMA-co-HEA)40/60, P(EMA-co-HEA)20/80 and PHEA) higher values of dc conductivity combined with strong temperature dependence, which was not quantitatively evaluated here, were observed. Obviously, PHEA becomes the continuous phase in P(EMA-co-HEA)40/60 (and compositions with higher PHEA content), whereas PEMA is the continuous phase in P(EMA-co-HEA)50/50 (and compositions with lower PHEA content) [58]. 12.3.5

Concluding Remarks

Results by broadband DS and TSDC on sequential IPNs with low and high crosslinking density and random copolymers of various poly(alkyl acrylate)s and poly(alkyl methacrylate)s presented in Section 12.3 illustrate the power of dielectric techniques for the investigation of mixing and phase separation in multiphase polymer systems. Several methodologies developed to that end have been introduced. TSDC thermograms provide a quick characterization of the overall dielectric behavior of the material under investigation (Figure 12.10).

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The main α relaxation associated with the glass transition can be followed by this technique even in polymeric systems with relatively high ionic conductivity (Figure 12.11), where DS fails. Significant information can be extracted already from the raw DS data with suitable choice of the formalism used and the quantity presented. Similar to TSDC, isochronal plots provide an overview of the overall dielectric behavior, however now the frequency of presentation can be selected in a wide range, stressing different aspects of that behavior (Figure 12.4). Isochronal ε (T) plots and Cole-Cole plots (Figure 12.7) provide directly quantitative information on dielectric strength, whereas reduced Cole-Cole plots (master curves) are directly conclusive with respect to the temperature dependence of the shape of a relaxation (Figure 12.9). Analysis of the DS data by fitting appropriate model functions (here the HN function (Eq. 12.6)), often a difficult task, provides detailed information on time scale, magnitude and shape of the response (Figures 12.5 and 12.6). Raw data and analysis provide several measures of Tg , such as T α by TSDC and T g diel by DS, which can be presented against composition facilitating further discussion on mixing and phase separation (Figures 12.8 and 12.12). Finally, dc conductivity provides significant information on phase morphology (phase continuity), as the moving ions probe the local morphology (Figure 12.13).

12.4 12.4.1

Rubber/Silica Nanocomposites (Interfacial Phenomena) Introduction

Polymer nanocomposites (PNCs), defined as composites where the filler is at least in one dimension smaller than about 100 nm, have attracted much interest in recent years for various technological applications, as well as for fundamental research. Several properties of the polymer matrix (such as mechanical, thermal and barrier properties) are significantly improved in PNCs at much lower filler factors, as compared to macroor microscale (conventional) composites [69, 70]. Despite many efforts [71, 72], there is as yet no complete theoretical explanation for that behavior. It is generally accepted, however, that interactions at the polymerfiller interfaces play a significant role. Results obtained by various techniques indicate the presence of an interfacial polymer layer around the filler, with structure/morphology and chain dynamics modified with respect to the bulk polymer matrix [73]. The existence of such an interfacial layer has been postulated for conventional composites long ago and various experiments provided support for that [74]. Questions related to the existence of such an interfacial layer, its thickness and the variation of polymer properties within the layer with respect to bulk properties become crucial for PNCs, as the interfacial layer can represent now a significant volume fraction of the polymer. Thus, PNCs become interesting also for fundamental studies of interfacial effects. A better understanding of these effects may provide a basis for understanding (and, thus, tailoring) the improvement of properties at the molecular level. We may expect that polymeric chains in the vicinity of a solid surface, within a distance of a few nm, exhibit different organization (density, chain conformation) and properties (thermal transitions, molecular mobility), as compared to chains in the bulk [75]. Computer simulations [76, 77] and experiments with model systems [78] provide support for that. Molecular dynamics simulations show that relaxation times may increase or decrease, as compared to the bulk material, depending on the type and strength of interaction and the roughness of the surface. A general result, obtained for a variety of materials and geometries, is that dynamics is similar to that of the bulk material far from the surface (distances larger than a few nm) and changes gradually and significantly (changes of a few orders of magnitude in relaxation times) by approaching the surface [76]. Scheidler et al. [77] proposed an empirical relationship for the dependence of relaxation time on the distance from the surface. Results on chain dynamics in PNCs, in particular segmental dynamics associated with the glass transition, reported in the literature often appear controversial and confusing [79]: dynamics (typically quantified in terms

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of glass transition temperature) may become faster or slower or show no change, it may be homogeneous or heterogeneous, etc. Obviously, several factors, such as polymer-filler interactions, filler size and morphology (degree of dispersion), may affect polymer dynamics in PNCs [80] and should be critically considered. Also, different experimental techniques, including mainly DSC to follow the glass transition, DMA, NMR and DS, each of them characterized by specific features, as indicated in Section 12.1, are typically employed in molecular dynamics studies in PNCs, a point which should be considered when discussing in terms of chain dynamics results obtained by various techniques. In this section we discuss results on polymer dynamics in rubber/silica nanocomposites. Experimental results in the literature obtained with rubber/silica and other PNCs have often been explained in terms of a three-layer model (strongly bound, loosely bound and quasi-bulk polymer), originally proposed by Tsagaropoulos and Eisenberg [81] and confirmed by other authors [82, 83]. Similar results have been explained also in terms of a two-layer model: a single interfacial layer with reduced dynamics and quasi-bulk polymer [84–86], or a continuous distribution of glass transition temperatures as a function of the distance from the particle surface [87]. Here we focus on poly(dimethylsiloxane) (PDMS)/silica NCs [88]. Details of the preparation of the unfilled, crosslinked PDMS and of the nanocomposites, by a sol–gel process in the presence of the crosslinked PDMS, have been given elsewhere [88]. The amount of filler incorporated into the network was varied between 0 and 16 vol%. In the following, samples are designated by PDMSX, where X is the volume fraction of silica. Transmission electron microscopy (TEM) images showed an excellent distribution of silica particles in the matrix with a diameter of about 10 nm. Next to morphology, polymer-filler interactions are also well characterized, as the particles interact with the matrix via hydrogen bonds between the oxygens on the PDMS backbone and the hydroxyls on the silica surface [88]. Thus, we are dealing with a system particularly suited to fundamental studies on interfacial effects on polymer dynamics. DSC, TSDC, DS, and DMA were employed to systematically study polymer dynamics in these PNCs [88– 90]. The DSC data were analyzed, next to crystallization not studied here [88], in terms of glass transition temperature Tg , heat capacity jump Cp at Tg , which is related to the fraction of polymer participating at the glass transition, and the width of transition, related to heterogeneity. Tg showed no significant variation with composition. However, the heat capacity jump at Tg , normalized to the fraction of amorphous polymer, Cp norm , systematically decreased with increasing filler content, indicating that a fraction of polymer does not participate to the glass transition being immobilized on the surface of the silica particles [88]. Assuming that the silica nanoparticles are spherical, have a diameter of 10 nm and are statistically distributed in the polymer matrix (Figure 12.14), the thickness of the immobilized layer was calculated to about 2 nm. The dielectric DS and TSDC data showed a different behavior, namely a slowing down of dynamics in the interfacial layer, as compared to bulk polymer, i.e. two segmental α relaxations. Finally, DMA data showed a moderate slowing down of segmental dynamics of the whole polymer matrix, manifested by an increase of the single Tg by about 10◦ C, as compared to the neat polymer [90]. In what follows we focus on aspects of the use of TSDC and DS in this investigation not stressed in Section 12.3. In particular, we make use of the high sensitivity and peak resolving power of TSDC and of the broad frequency range of DS in combination with sophisticated methods of data treatment to analyze and quantify polymer dynamics in the interfacial layer [88–90]. 12.4.2

TSDC Studies

Figure 12.15 shows TSDC thermograms recorded with neat PDMS and four NCs. The thermogram in neat PDMS shows a single peak at about 150 K, which corresponds to the segmental α relaxation associated with the glass transition of the amorphous phase of PDMS, in good agreement with the DSC data [88]. For the NCs the α relaxation is observed at approximately the same temperature, but with a higher intensity due to the concomitant decrease of the degree of crystallinity [88]. In addition, a shoulder is observed on the high-temperature side of the main peak extending up to approximately 30 K higher. The temperature

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Figure 12.14 Model depicting the morphology of the PDMS/silica NCs. Reprinted from [88]. Copyright (2005) with permission from Elsevier.

position of the shoulder is independent of composition, whereas its intensity increases systematically with silica content. The shoulder is assigned to the α relaxation of PDMS chains in an interfacial layer close to the silica particles (in the following α  relaxation), where chain mobility is constrained due to interaction with the surface of the particles (hydrogen bonding). Bearing in mind Eq. (12.14), the relative areas of the main relaxation and the shoulder give directly the relative volume of the interfacial and bulk phases. For direct comparison the thermograms in the region of the glass transition have been normalized to the same height of the main peak (Figure 12.16). The contribution of the shoulder (α  relaxation, dotted line) is then obtained by subtracting the pure PDMS signal from the normalized NCs signal. The dielectric strength ε of the α relaxation and of the α  relaxation is then calculated. The results, listed in Table 12.1, show that, with increasing silica content, the dielectric strength of the α relaxation decreases in general, whereas that of the α  relaxation increases approximately linearly. Using the model of interfacial layer of Figure 12.14, the fraction of polymer with reduced mobility and the thickness of the interfacial layer are calculated (Table 12.1). Interestingly, values similar to those calculated from the DSC data are obtained.

Figure 12.15 TSDC thermograms for pure PDMS and the PDMS/silica NCs indicated on the plot. Reprinted from [88]. Copyright (2005) with permission from Elsevier.

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Figure 12.16 TSDC thermograms in the region of the glass transition in pure PDMS and two NCs normalized to the temperature and height of the main peak. The dashed line, obtained by subtracting the PDMS thermogram from the nanocomposite thermogram, gives the contribution of the α  relaxation. Reprinted from [88]. Copyright (2005) with permission from Elsevier.

At temperatures higher than those of the α and the α  TSDC peaks a sharp peak is observed in the thermograms of Figure 12.15 for all the samples at about 225 K and a broad peak at 250–290 K. The sharp peak is associated with melting in the crystalline regions of PDMS, in agreement with DSC results [88], illustrating the power of TSDC for first order phase transition studies. The broad peak at 250–290 K, which increases in magnitude and shifts to higher temperatures with increasing silica content, is assigned to interfacial Maxwell-Wagner-Sillars (MWS) polarization/relaxation, i.e. to trapping of charge carriers (ions) at interfaces between regions of different conductivity during the polarization step and their release during the depolarization step [50, 51]. Two kinds of interfaces coexist in the PDMS/silica interfaces at low temperatures, those between crystalline and amorphous regions and those between polymer and filler. Since the MWS peak is recorded at temperatures higher than the melting temperature, it must be totally assigned to polymer/filler interfaces. The magnitude of the peak increases significantly for the NC with the largest (16%) silica content, which might indicate a qualitative change of the morphology for that composition, more likely the existence of polymer regions surrounded by silica nanoparticles resulting in an enhancement of trapping of charge carriers. Thus, detailed investigation of the MWS peak by TSDC may provide significant information on morphology [5, 9]. The TSDC results in Figures 12.15 and 12.16 are consistent with both (i) two distinct α and α  relaxations associated with two distinct glass transitions, and (ii) a continuous distribution of relaxation times and glass Table 12.1 Normalized dielectric strengths ε α and εα of the α and α  relaxations, respectively, fraction χ int of interfacial polymer obtained from these and corresponding thickness dint of the interfacial layer. Sample

εα

εα

χ int

dint [nm]

PDMS0 PDMS6 PDMS9 PDMS10 PDMS16

1.12 1.13 0.88 0.96 1.00

– 0.36 0.46 0.55 0.80

– 0.13 0.21 0.22 0.29

– 2.3 2.5 2.5 2.2

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Figure 12.17 Thermal sampling responses in the region of the glass transition in pure PDMS and in a nanocomposite. The arrows indicate the polarization temperature for each response. Reprinted from [88]. Copyright (2005) with permission from Elsevier.

transition temperatures in the interfacial layer. TSDC offers special methods for experimentally analyzing complex relaxations into approximately single responses: variation of polarization temperature within the region of the relaxation, partial heating and, in particular, thermal sampling (TS) [50, 51]. Here the TS technique was employed. It consists of ‘sampling’ the relaxation processes within a narrow temperature range by polarizing at a temperature Tp and depolarizing at Td , a few degrees lower than Tp . The sample is then cooled down and the depolarization current measured during heating as in standard TSDC. In a series of TS measurements the polarization temperature is varied (here in steps of 5 K) to span the whole temperature range of the complex peak. As an example, Figure 12.17 shows TS responses (approximately single responses) experimentally isolated in the temperature region of the glass transition in pure PDMS (a), and in the nanocomposites with 10 vol% silica (b). In pure PDMS the TSDC α peak is characterized by a distribution of relaxation times, as commonly found in polymers [53, 91]. In the NCs the distribution becomes much broader at higher temperatures (longer relaxation times), corresponding to the region of the shoulder in Figure 12.15. Moreover, the magnitude (peak height) of the TS responses decreases continually with increasing polarization temperature and no sign is observed for a second TSDC peak in the temperature region of the shoulder. These results provide strong experimental evidence for a continuous distribution of relaxation times in the glass transition region of the NCs, corresponding to a continuous distribution of Tg ’s between that of pure PDMS and approximately 190 K [90]. The apparent activation energy W of the individual TS responses was calculated by the initial rise method, Eq. (12.12), and an approximate expression based on the shape of the peak (compare Section 12.2.3). The results clearly indicate the presence of cooperative phenomena in the region of both the α and α  relaxations [90]. 12.4.3

Broadband DS Studies

The dynamic glass transition, α and α  relaxations in pure PDMS and the NCs, were studied in detail by broadband DS [88–90]. Typical loss spectra are shown in Figure 12.18. Two loss peaks are visible at each temperature, assigned, in the order of increasing frequency, to the segmental relaxation of the interfacial layer and of the bulk PDMS, α  and α, respectively [89]. The data were analyzed in terms of time scale and relaxation strength, f max and ε, respectively by fitting a sum of HN terms to the data, Eq. (12.6), an example of the quality of fitting being given in Figure 12.18. A sum of three relaxations was necessary to reproduce the shape of the spectra, the third, weak relaxation being assigned to conductivity on the surface of

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Figure 12.18 Dielectric loss against frequency for the PDMS/silica NC containing 10% silica. Temperatures are from 158 to 183 K in steps of 5 K. The bold line is a fit of a sum of 3 HN terms, Eq. (12.6), to the data for 173 K, the thinner lines are the individual HN terms.

the silica particles due to adsorbed water molecules [88]. The spectra for the other NCs are similar to those in Figure 12.18, differing only in the relative magnitudes of the three relaxations. Figure 12.19 shows the Arrhenius plot for pure PDMS and a NC, with TSDC and DMA data included. Please note that for pure PDMS the DMA point is shifted by about 10 K to higher temperatures as compared to the dielectric data, which can be discussed in terms of a different spatial scale of dielectric and mechanical techniques [9]. At high temperatures the α and α  relaxations are well separated, their relaxation times differing by 3–4 decades. However, the interfacial relaxation has a weaker temperature dependence than the

Figure 12.19 Arrhenius plot of pure PDMS (open symbols) and the NC with 10 vol% silica (filled symbols). Included also are TSDC data as a point for α (main peak) and as a horizontal bar for α  (the shoulder) at the equivalent frequency of 1.6 mHz, and DMA E data for pure PDMS (x) and the nanocomposite (+) at 1 Hz. The lines are fits to the DRS data of the Arrhenius equation (12.16) for the intermediate relaxation and of the Vogel-Tammann-Fulcher equation (12.15) to the two α relaxations. Reprinted from [88]. Copyright (2005) with permission from Elsevier.

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Figure 12.20 Dielectric loss ε against temperature for pure PDMS and two NCs at three frequencies indicated on the plot. Reprinted from [88]. Copyright (2005) with permission from Elsevier.

bulk α relaxation and a smaller curvature, having an almost Arrhenius behavior. Thus, at lower temperatures and lower frequencies, approaching the glass transition, the α  relaxation tends to converge with the bulk α. This behavior is in agreement with the TSDC results, which correspond to lower frequencies (10−3 Hz) than those accessible with DS, and where the α  relaxation appears as a shoulder on the α peak. To make clearer the last point mentioned in the previous paragraph and to compare with measurements in the temperature domain (DSC, TSDC and DMA) we show in Figure 12.20 DS results (dielectric loss ε ) measured isothermally in frequency scans and replotted as a function of temperature for pure PDMS and two NCs at three frequencies. A single loss peak is observed in pure PDMS, in agreement with the DS data in the frequency domain [88] shifting to higher temperatures with increasing frequency. In the NCs a double peak is observed, the α loss peak being located at the same temperature as in pure PDMS and the α  loss peak at higher temperatures, becoming more separated and distinguished with increasing frequency. The overall increase of the response of the NCs in Figure 12.20 is related partly to the reduction of the degree of crystallinity [88] and partly to the increase of the internal electric field (compare Section 12.5.2). We now turn our attention to the second source of information extracted by analysis of the DS data, relaxation strength ε of the α and α  relaxations, presented in Figure 12.21 for two NCs [89]. In the temperature range of our data ε of the α increases slightly with increasing temperature and that of the α  process decreases, while the sum of the two remains approximately constant. By making the simplification that we are dealing with two types of polymer chains, bulk-like and interfacial, the relative dielectric strength of the two relaxations is a measure of the relative amount of bulk-like and interfacial polymer. Making the additional approximation that the interfacial polymer does not crystallize and using the model of Figure 12.14, we can estimate the fraction of interfacial polymer and from that the thickness d of the interfacial layer, similar

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1 α 0.8 dielectric strength Δε

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0.4

0.2

0

α′

160

170

180 190 temperature [K]

200

210

Figure 12.21 Dielectric strength against temperature for the α (circles) and α’ (squares) relaxations, for the PDMS/silica NCs with 6% (open symbols) and 9% silica (filled symbols).

to what we did and in good agreement with the DSC and the TSDC data (Table 12.1), now, however, as a function of temperature. One can argue that d is related to the cooperativity length of the glass transition ξ [11, 16] and that the main factor that influences d is in fact cooperativity [89]. The temperature dependence of d can then be discussed in relation to models for the temperature dependence of ξ , as well as to results of molecular dynamics simulations [89]. The interplay between ξ and d can also explain the apparent inconsistency between the TSDC results, where there is no well-defined second Tg but a continuous distribution, and the DS results where a distinct second α relaxation several decades slower is observed: at the glass transition these two characteristic lengths are comparable, however, with increasing temperature d is found to be relatively constant but the ξ decreases significantly, allowing the appearance of a second distinct relaxation. A final comment in this section refers to effects of interfacial interactions in PNCs on dielectric relaxations with characteristic lengthscales smaller (local, secondary relaxations) and larger (global, normal mode relaxation) than that of the segmental α relaxation. Results show that in the first case effects, in particular on the time scale of the relaxation, are rather negligible [64, 92], which can be understood in terms of the local character of these relaxations. An example will be given in Section 12.5. Results are less clear for the normal mode relaxation (global relaxation) [93], which is related with the presence of a dipole moment component along the chain contour [12], as it is usually overlapped with and often masked by dc conductivity and related effects [94]. 12.4.4

Concluding Remarks

Results on polymer dynamics in PDMS/silica NCs presented in Section 12.4 demonstrate the power of dielectric techniques for the investigation of interfacial effects in PNCs. TSDC thermograms enable to follow, next to dipolar relaxations, first order transitions such as melting, whereas the investigation of the interfacial MWS relaxation provides significant information on morphology (Figure 12.15). Complex relaxations can

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be experimentally analyzed into approximately single responses by the thermal sampling technique, a special variant of TSDC, and the results allow to distinguish between a continuous distribution of relaxation times and overlapping of discrete peaks (Figure 12.17). The broad frequency range of DS combined with sophisticated methods of critical analysis of the experimental data (Figure 12.18) enable to study in detail the time scale of the bulk and the interfacial segmental relaxations (Figure 12.19) and the temperature dependence of their strengths (Figure 12.21), which are interpreted in terms of fractions of bulk and interfacial polymer. The temperature dependence of the thickness of the interfacial layer d can then be discussed in relation to models for the temperature dependence of the cooperativity length of the glass transition ξ and to results of molecular dynamics simulations. Critical comparison of dielectric (DS and TSDC), DSC and DMA results obtained with the same well-characterized system reveals distinct similarities and differences, which can be discussed in terms of time and spatial scale of each technique (Section 12.4.1).

12.5

12.5.1

Polymer Nanocomposites with Conductive Carbon Inclusions (Percolation Phenomena) Introduction

Polymer NCs with conductive inclusions (in the following, conductive PNCs) have attracted much interest in recent years for various applications, such as electromagnetic shielding coatings, electrostatically dissipative materials, aerospace structural materials and active elements in sensors [95]. This class of materials is a typical example of composites, where we utilize properties of the inclusion, namely electrical conductivity, whereas the matrix (a polymer) provides other good properties depending on the specific application envisaged, such as processability and mechanical stability. Thus, a main issue under investigation in these materials is electrical conductivity, in particular its dependence on filler content, size, shape and morphology of dispersion, i.e. percolation [96]. The quality of conductive PNCs is often evaluated in terms of the percolation threshold (pc ), which is the critical concentration of the filler where conducting pathways are formed by the nanoparticles and consequently a transition from the insulating to the conducting phase is observed. pc is usually determined and percolation is studied through dc conductivity measurements. However, much more information on electrical conductivity and percolation can be obtained by ac conductivity measurements [48, 49, 97, 98], i.e. broadband DS (compare Section 12.2.2). The main task in this section is to illustrate that point. We will demonstrate also the suitability of ac conductivity measurements for evaluating the quality of filler dispersion and, thus, the power of the technique for monitoring and controlling that dispersion. The control of filler dispersion is a central task in the science and technology of PNCs [99]. Other aspects of the use of dielectric spectroscopy to be discussed in this section refer to (i) the investigation of the frequency dependence of the real part of the dielectric function, ε (f ), with respect to the determination of the percolation threshold pc , as well as to the search for materials with high values of ε (high k materials), and (ii) the investigation of effects of the filler on polymer dynamics (interfacial phenomena), in particular secondary relaxations. Examples illustrating the points mentioned above will be taken from two families of PNCs, an epoxy resin system with conducting spherical carbon nanoparticles and semicrystalline thermoplastics, polypropylene and polyamide 6, with carbon nanotubes. 12.5.2

Analysis of DS Data in Terms of the Dielectric Function

Figure 12.22 shows results for the real and imaginary part of the dielectric function recorded by DS on PNCs prepared from araldite (bisphenol A epoxy resin), diethylenetriamine as hardener and nanosized carbon particles (NCPs), a mixture of graphite (∼3 nm) and diamond (∼10 nm) particles (in a ratio of 67:33) as filler

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20 ε' 16

16

12

1.2

8

14 12

1.4

2

4

6

carbon vol%

0.8

10 8 6

ER 1% NCP 4% NCP 6% NCP 80805 Hz

1.0

4 0

ε''

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0.4 0.2

4 -160 -120 -80 -40

0.6

0

40

80 120 160

0.0

-160 -120 -80 -40

0

40 80 120 160

o

o

T [ C]

T [ C]

(b)

(a)

Figure 12.22 Real part (ε ) (a) and imaginary part (ε ) (b) of the dielectric function versus temperature (T) of the ER/NCP NCs indicated on the plot at about 80 kHz. The inset shows ε  (measured at 1 Hz and −50◦ C) versus the NCP volume concentration. The line is a fit of Eq. (12.18) to the data. Reprinted from [100]. Copyright (2005) with permission from John Wiley & Sons.

[100]. The data were recorded isothermally by scanning the frequency and replotted here as a function of temperature (isochronal plot). A relatively high frequency (about 80 kHz) has been chosen for the presentation, in order to eliminate conductivity effects present at lower frequencies [100]. An overall increase of molecular mobility is observed in Figure 12.22, in agreement with TSDC data not shown here, in the sense that, at each temperature, ε and ε increase with increasing filler content. This is to a large extent due to an increase of the internal field [94], related with the formation of a percolation structure of the nanoparticles. We will comment on implications from such measurements on the evaluation of dielectrics based on PNCs with conductive inclusions as materials with high values of ε (high k materials), e.g. for capacitor applications, later in this section. The inset to Figure 12.22(a) shows the dependence of ε (at a frequency of 1 Hz and a temperature of –50◦ C) on volume concentration of carbon. We have used that dependence to calculate the percolation threshold pc , by fitting the following equation from percolation theory [96, 100] to the experimental ε (p) data: ε ( p) = εm + A | p − pc |−t

(12.18)

where εm is the dielectric constant of the matrix, pc the percolation threshold and t the critical exponent. From the results shown in Figure 12.22(a), pc and t have been determined to 7.4% and 0.69, respectively. Two comments are in order here from the methodological point of view with respect to the use of Eq. (12.18) for percolation studies. The first is that this equation provides a second, independent possibility to study percolation and determine the percolation threshold pc by means of DS, not often used, in addition to the well-known one based on conductivity, to be discussed in the next section. The second comment is that information on percolation obtained from Eq. (12.18) (and from conductivity measurements) refer to electrical percolation and can be critically compared with similar information obtained by other techniques, such as rheological measurements on the system taken here as example [100]. Figure 12.23 shows DS results for ε (f ) at room temperature for PNCs based on polypropylene grafted with maleic anhydride (MA-PP) and multi-walled carbon nanotubes (CNTs) [101]. They refer to compositions below the percolation threshold, determined by means of Eq. (12.18) from the results shown in the figure and, independently, by analysis of ac conductivity data (compare next section) to 2.2 wt% [101]. We observe in

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Figure 12.23 Real part ε of the dielectric function versus frequency f at room temperature (25◦ C) for the MA-PP/CNT NCs indicated on the plot.

Figure 12.23 that ε increases with increasing filler content, the additional contribution obviously arising from interfacial Maxwell-Wagner-Sillars (MWS) polarization (compare Section 12.4.2) of isolated or aggregates of CNTs. ε in Figure 12.23 is practically independent of frequency for all compositions, also at other temperatures not shown here, for two reasons. The first is that PP is a nonpolar polymer. Maleic anhydride is a polar group, which is attached to the PP backbone; nevertheless, its contribution to ε is negligible, as indicated by the results in the figure for the neat matrix, because of its very low content (0.14 wt%). The second is that the high conductivity of CNTs shifts the MWS relaxation to higher frequencies in the GHz region [13], so that the additional MWS polarization is also practically frequency-independent. These two features, increase of ε combined with frequency- and temperature-independence, render these materials interesting for capacitor applications [102]. Two relaxations, a secondary β relaxation at lower temperatures and the primary α relaxation at higher temperatures, associated with the glass transition of the epoxy resin matrix, are observed in Figure 12.22. For both relaxations the strength, i.e. the magnitude of the peak ε (T), and the corresponding step in ε (T) increase in the nanocomposites, in particular for the sample above the percolation threshold. The time scale (temperature position) of the response shows, however, a different behavior. For the local β relaxation it does not change with composition, whereas for the cooperative α relaxation the peak temperature increases slightly in the nanocomposites, in particular at higher filler contents, where it shifts out of the temperature range of Figure 12.22. Measurements at higher temperatures/lower frequencies are less conclusive for the higher content nanocomposites, as the results are masked by conductivity effects. Thus, similar to the rubber/silica NCs in Section 12.4 for the α relaxation, the results of DS provide information on molecular mobility in terms of the relaxation strength and the time scale of the response, now also for the local β relaxation. The increase of relaxation strength for both β and α can be understood in terms of increased free volume, in agreement with results obtained with other nanocomposites [64, 92], and/or increase of the internal field [94]. The slowing down of the cooperative α relaxation (dynamic glass transition), which corresponds to an increase of Tg ,

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may be understood in terms of immobilization of polymer chains in the interface layer around the particles (formation of bound polymer), similar to the case of rubber/silica NCs in Section 12.4, or an effect of NCPs on curing (increase of crosslinking density), as polymerization and curing was performed in the presence of the filler [92]. The timescale of the β relaxation, on the contrary, is less affected by both interfacial interactions and increase of crosslinking density, because of the local character (shorter lengthscale) of the relaxation.

12.5.3

Analysis of DS Data in Terms of ac Conductivity

Figure 12.24 shows results for the frequency dependence of ac conductivity, σ ac (f), of PNCs based on poly(amide 6) (PA) and multi-walled CNTs in a wide range of compositions at 25◦ C. Details of preparation of the materials and of characterization by various techniques have been given elsewhere [103]. σ ac (f ) has been calculated by Eq. (12.8) from the measured ε (f ) data. For neat PA and PA/2.5 wt% CNT ac conductivity in Figure 12.24 increases approximately linearly with increasing frequency in logarithmic scale, exhibiting a typical capacitor behavior. At loadings in excess of 5 wt% a dc plateau, where conductivity is independent of frequency, appears up to the critical frequency fc . The crossover frequency fc separates the region of macroscopic charge carrier transport at lower frequencies and their confined motion within their potential wells at higher frequencies. For these composites fc increases with increasing CNT content and shifts out of the frequency range of measurements for the sample PA/20 wt% CNT. Thus, the dc conductivity plateau is clearly achieved above 2.5 and below 5.0 wt% CNT, indicating that the percolation threshold pc , the transition from the insulating to the conducting phase, is located in the range 2.5–5.0 wt% (or 1.6–3.3 vol%) CNT. Analysis of the results for the conducting samples in Figure 12.24 in terms of the universal dielectric response, Eq. (12.9), or other similar expressions provides information on the mechanism of charge carrier transport [98, 103]. Figure 12.25 shows results for the dependence of dc conductivity σ dc , determined by the plateau values in Figure 12.24, on CNT volume content p for the PA/CNT NCs above the percolation threshold. The well-known

Figure 12.24

Frequency dependence of ac conductivity for the PA/CNT NCs indicated on the plot at 25◦ C.

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Figure 12.25 σ dc versus CNT volume concentration for PA/CNT NCs above pc . The inset shows a log-log plot of σ dc versus (p-pc ), Eq. (12.19) with t = 8.4 and pc = 1.7 vol. %.

scaling law [97–99]: σdc ( p) ∼ ( p − pc )t

(12.19)

where t is a critical exponent related with the dimensionality of the system under investigation, was fitted to the experimental data and the values of pc and t were determined to 1.7 vol% and 8.4, respectively. The significance of these results has been discussed in [103]. The results of detailed studies indicate that electrical conductivity in polymer/CNT NCs depend on the type of polymer and CNT used and the degree of dispersion of the latter, i.e. the processing conditions [95, 98, 101–104]. The dependence on the degree of dispersion is strong, also for PNCs with other conducting inclusions [92, 100], suggesting that monitoring of conductivity might be a powerful technique for evaluating the effects of various processing conditions on the quality of filler dispersion. This can be done either at the final stage of sample preparation, i.e. on solid samples pressed or cast from solution, or during mixing in the melt/solution. As explained in Section 12.5.1, the use of ac conductivity measurements provides certain advantages over that of dc conductivity measurements. Figure 12.26 shows results of ac conductivity measurements at room temperature on PNCs based on poly(propylene) and multiwalled CNTs at a fixed content of 5 wt% filler and different mixing conditions in the melt: mixing speed, temperature and time of mixing. After mixing in a Plasti-Corder kneading machine, solid samples for ac conductivity measurements were prepared by compression molding using a laboratory hydraulic press. Details of these measurements will be published elsewhere. The results in Figure 12.26 show that the overall behavior, in particular dc conductivity, determined by the plateau values, and crossover frequency fc , depend sensitively on mixing conditions and their measurement can be used to evaluate the quality of filler dispersion.

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Figure 12.26 Frequency dependence of ac conductivity at 25◦ C for PP/CNT NCs with 5 wt% MWCNTs and the conditions of mixing indicated on the plot.

12.5.4

Concluding Remarks

Results on PNCs with conducting spherical carbon nanoparticles and carbon nanotubes presented in Section 12.5 demonstrate the power of ac conductivity measurements (Figure 12.24) for a detailed investigation of percolation and charge transfer mechanism in these systems. Next to accurate values of dc conductivity, used for the determination of percolation thrshold and critical exponent (Figure 12.25), ac conductivity plots provide additional information on percolation and charge transfer through the crossover frequency from dc to ac conductivity and the frequency dependence of ac conductivity (Eq. (12.9)). Ac conductivity measurements are particularly suited for evaluating the quality of filler dispersion (Figure 12.26), rendering DS a powerful tool for monitoring and controlling that dispersion. The investigation of the frequency dependence of the real part of the dielectric function, ε (f ), provides an independent way for quantitative characterization of percolation (Figure 12.22), next to the evaluation of dielectric materials as high k materials (Figure 12.23).

12.6 Conclusion Dielectric techniques, including broadband DS and TSDC, are a powerful tool in the investigation of polymer dynamics in multiphase polymer systems. Evidence for that is provided at three selected examples presented in this chapter: random copolymers and interpenetrating polymer networks of poly(alkyl acrylate)s and poly(alkyl methacrylate)s, where we focused on polymer dynamics in relation to mixing and microphase separation; rubber/silica nanocomposites with a fine dispersion of silica nanoparticles, where we focused on the investigation of effects of interfacial interactions between the two components on polymer dynamics; and polymer nanocomposites with conducting carbon inclusions, where we demonstrated the power of dielectric techniques to study electrical conductivity in such systems, percolation and charge transfer between conducting islands.

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TSDC provides a quick characterization of the overall dielectric behavior of the system under investigation. The method is characterized by simplicity, high sensitivity and high resolving power and it provides special variants, in particular the thermal sampling technique, for experimentally analyzing complex relaxations into approximately single responses. It enables to follow the segmental α relaxation associated with the glass transition in systems with relatively high ionic conductivity, where DS fails, and it is particularly suited for the investigation of interfacial polarization/relaxation, providing significant information on morphology. The main feature of DS, which renders it attractive for detailed studies of dynamics and of structure/morphology in multiphase polymer systems, is the extremely broad frequency range covered by the technique. This enables to follow on the same sample and under the same conditions processes with very different time scales, from shorter than picoseconds to longer than hours, corresponding to spatial scales from tenths of a nanometer to several micrometers, the latter in relation to conductivity studies. The broad frequency and temperature ranges of measurements, combined with the use of various formalisms of data presentation, which stress different aspects of molecular dynamics and charge carrier motion, enable to extract significant quantitative information already from the raw data, as illustrated at several places in this chapter. Critical analysis of the data by fitting appropriate model functions, often by sophisticated methodologies, is however necessary for gaining full quantitative information on time scale, strength and shape of the response. Results of dielectric studies on polymer dynamics in multiphase polymer systems are typically discussed in relation to those obtained by other techniques, in particular DSC and DMA in the case of the three selected examples of this chapter. Critical comparison of the results by various techniques reveals distinct similarities and differences, which can be discussed in terms of time and spatial scale of each technique. We would like to stress here the need for a complete analysis of the whole set of data provided by each technique for a correct interpretation of the experimental results, which include, for example, heat capacity jump at Tg , next to Tg , in the case of DSC, and strength and shape of a relaxation, next to time scale, in the case of DS.

Acknowledgements The authors wish to thank their colleagues J. L. Gomez Ribelles, M. Monleon Pradas, G. Gallego Ferrer, J. M. Meseguer Duenas, S. Kripotou, L. Bokobza and R. Kotsilkova for providing data and/or for valuable discussions.

References 1. M. Alexandre and P. Dubois, Polymer-layered silicate nanocomposites: preparation, properties and use of a new class of materials, Mater. Sci. Eng., 28, 1–63 (2000). 2. S. Sanchez, B. Lebeau, F. Chaput, and J.-P. Boilot, Optical properties of functional hybrid organic-inorganic nanocomposites, Adv. Mater., 15, 1969–1994 (2003). 3. S. S. Ray and M. Okamoto, Polymer/layered silicate nanocomposites: a review from preparation to processing, Prog. Polym. Sci., 28, 1539–1641 (2003). 4. S. H. Phillips, T. S. Haddad, and S. J. Tomczak, Developments in nanoscience: polyhedral oligomeric silsequioxane (POSS)-polymers, Curr. Opin. Solid State Mater. Sci, 8, 21–29 (2004). 5. P. Pissis, A. Kyritsis, G. Georgoussis, and V. Shilov, Structure-property relationships in polyurethane ionomers, in Advanced Functional Molecules and Polymers Vol. 1, H. S. Nalwa (Ed.), Gordon and Breach Publishers, Amsterdam, 2001. 6. V. A. Bershtein and V. M. Egorov, Differential Scanning Calorimetry of Polymers. Physics, Chemistry, Analysis, Technology, Ellis Horwood, Chichester (1994). 7. D. Campbell and J. R. White, Polymer Characterization, Chapman and Hall, London, 1989.

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8. S. J. Spell (Ed.), Characterization of Solid Polymers, Chapman and Hall, London, 1994. 9. A. S. Vatalis, A. Kanapitsas, C. G. Delidis, and P. Pissis, Relaxation phenomena and morphology in polymer blends based on polyurethanes investigated by various thermal analysis techniques, Thermochim. Acta, 372, 33–38 (2001). 10. M. Doxastakis, M. Kitsiou, G. Fytas, D. N. Theodorou, N. Hadjichristidis, G. Meier, and B. Frick, Component segmental mobilities in an athermal polymer blend: Quasielastic incoherent neutron scattering versus simulation, J. Chem. Phys., 112, 8687–8694 (2000). 11. E. Donth, The Glass Transition. Relaxation Dynamics in Liquids and Disordered Materials, Springer, Berlin, 2001. 12. A. Kyritsis, P. Pissis, S.-M. Mai, and C. Booth, Comparative dielectric studies of segmental and normal mode dynamics in poly(oxybutylene) and poly(oxyethylene)-poly(oxybutylene) diblock copolymers, Macromolecules, 33, 4581–4595 (2000). 13. P. Hedvig, Dielectric Spectroscopy of Polymers, Adam Hilger, Bristol, 1977. 14. J. M. G. Cowie, S. Harris, J. L. Gomez Ribelles, J. M. Meseguer, F. Romero, and C. Torregrosa, Glass transition and structural relaxation in polystyrene/poly(2,6-dimethyl-1,4-phenylene oxide)miscible blends. Macromolecules, 32, 4430–4438 (1999). 15. S. K. Kumar, R. H. Colby, S. H. Anastasiadis, and G. Fytas, Concentration fluctuation induced dynamic heterogeneities in polymer blends, J. Chem. Phys., 105, 3777–3788 (1996). 16. G. Adam and J. H. Gibbs, On the temperature dependence of cooperative relaxation properties in glass-forming liquids, J. Chem. Phys., 43, 139–146 (1965). 17. P. Maroulas, S. Kripotou, P. Sysel, R. Hobzova, J. Kotek, and P. Pissis, Molecular dynamics in hyperbranched polyimides cross-linked with ethylene glycol diglycidyl ether, J. Non-Cryst. Solids, 352, 4800–4803 (2006). 18. S. Kamath, R. H. Colby, S. K. Kumar, K. Karatasos, G. Floudas, G. Fytas, J. E. L. Roovers, Segmental dynamics of miscible polymer blends: comparison of the predictions of a concentration fluctuation model to experiment, J. Chem. Phys., 111, 6121–6128 (1999). 19. Y. H. Chin, C. Zhang, P. Wang, P. T. Inglefield, A. Jones, R. P. Kambour, J. T. Bendler, and D. M. White, Glass transition dynamics in a compatible blend by two-dimensional solid-state NMR, Macromolecules, 25, 3031–3038 (1992). 20. G. Katana, E. W. Fischer, T. Hack, V. Abetz, and F. Kremer, Influence of concentration fluctuastions on the dielectric α-relaxation in homogeneous polymer mixtures, Macromolecules, 28, 2714–2722 (1995). 21. G.-C. Chung, J. A. Kornfield, and S. D. Smith, Component dynamics miscible polymer blends: a two-dimensional deuteron NMR investigation, Macromolecules, 27, 964–973 (1994). 22. T. P. Lodge and T. C. B. McLeish, Self-concentration and effective glass transition temperatures in polymer blends, Macromolecules, 33, 5278–5284 (2000). 23. T. P. Lodge, E. R. Wood, and J. C. Haley, Two calormetric glass transitions do not necessarily indicate immiscibility: The case of PEO/PMMA, J. Polym. Sci. Part B Polym. Phys., 44, 756–763 (2006). 24. D. Cangialosi, G. A. Schwartz, A. Alegria, and J. Colmenero, Combining configurational entropy and selfconcentration to describe the component dynamics in miscible polymer blends, J. Chem. Phys., 123, 144908 (2005). 25. G. A. Schwartz, D. Cangialosi, A. Alegria, and J. Colmenero, Describing the component dynamics in miscible polymer blends: Towards a fully predictive model, J. Chem. Phys., 124, 154904 (2006). 26. I. Cendoya, A. Alegria, J. M. Alberdi, J. Colmenero, H. Grimm, D. Richter, and B. Frick, Effect of blending on the PVME dynamics. A dielectric, NMR, and QENS investigation, Macromolecules, 32, 4065–4078 (1999). 27. J. W. Sy and J. Mijovic, Reorientational dynamics of poly(vinylidene fluoride)/poly(methyl methacrylate) blends by broad-band dielectric relaxation spectroscopy, Macromolecules, 33, 933–946 (2000). 28. M. Dionisio, A. C. Fernandes, J. F. Mano, N. T. Correia, and R. C. Sousa, Relaxation studies in PEO/PMMA blends, Macromolecules, 33, 1002–1011 (2000). 29. A. Alegria, D. Gomez, and J. Colmenero, Temperature-pressure eqvivalence for the component segmental dynamics of a miscible polymer blend, Macromolecules, 35, 2030–2035 (2002). 30. G. Floudas, G. Fytas, T. Reisinger, and G. Wegner, Pressure-induced dynamic homogeneity in an athermal diblock copolymer melt, J. Chem. Phys., 111, 9129–9132 (1999). 31. C. M. Roland, K. J. McGrath, and R. Casalini, Dynamic heterogeneity in poly(vinyl methyl ether)/poly(2chlorostyrene) blends, Macromolecules, 39, 13581–3587 (2006).

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32. Y. Hirose, O. Urakawa, and K. Adachi, Dynamics in disordered block copolymers and miscible blends composed of poly(vinyl ethylene) and polyisoprene, J. Polym. Sci. Part B Polym. Phys., 42, 4084–4094 (2004). 33. N. Miura, W. J. MacKnight, S. Matsuoka, and F. E. Karasz, Comparison of polymer blends and copolymers by broadband dielectric analysis, Polymer, 42, 6129–6140 (2001). 34. X. Jin, S. Zhang, and J. Runt, Dielectric study of blends of poly(ethylene oxide) and poly(styrene-co-phydroxystyrene). Semicrystalline blends, Macromolecules, 37, 4808–4814 (2004). 35. J. Ren and K. Adachi, Dielectric relaxation in blends of amorphous poly(dl-lactic acid) and semicrystalline poly(llactic acid), Macromolecules, 36, 5180–5186 (2003). 36. J. C. Haley, T. P. Lodge, Y. He, M. D. Ediger, E. D. von Meerwall, and J. Mijovic, Composition and temperature dependence of terminal and segmental dynamics in polyisoprene/poly(vinylethylene) blends, Macromolecules, 36, 6142–6151 (2003). 37. S. Zhang, P. C. Painter, and J. Runt, Suppression of the dielectric secondary relaxation of poly(2-vinylpyridene) by strong intermolecular hydrogen bonding, Macromolecules, 37, 2636–2642 (2004). 38. V. N. Daniel, Dielectric Relaxation, Academic Press, London, 1967. 39. N. G. McGrum, B. E. Read, and G. Williams, Anelastic and Dielectric Effects in Polymeric Solids, Wiley, New York, 1967. 40. J. Runt and J. J. Fitzgerald (Eds.), Dielectric Spectroscopy of Polymeric Materials, American Chemical Society, Washington, DC, 1997. 41. F. Kremer and A. Schoenhals (Eds.), Broadband Dielectric Spectroscopy, Springer, Berlin, 2003. 42. S. Takeishi and S. Mashimo, Dielectric relaxation measurements in the ultralow frequency range, Rev. Sci. Instrum., 53, 1155–1159 (1982). 43. S. Nozaki and T. K. Bose, Broadband complex permittivity measurements by time-domain spectroscopy, IEEE Trans. Instrum. Meas., 39, 945–951 (1990). 44. Y. Feldman, A. Andrianov, E. Polygalov, I. Ermolina, G. Romanychev, Y. Zuev and B. Milgotin, Time domain dielectric spectroscopy: an advanced measuring system, Rev. Sci. Instrum., 67, 3208–3216 (1996). 45. S. Havriliak, Jr and S. J. Havriliak, Dielectric and Mechanical Relaxation in Materials, Hanser Verlag, Munich, 1997. 46. C. T. Moynihan, L. P. Boesch and N. L. Laberge, Decay function for the electric field relaxation in vitreous ionic conductors, Physics Chem. Glasses, 14, 122–125 (1973). 47. J. R. Macdonald (Ed.), Impedance Spectroscopy, John Wiley & Sons Inc., New York, 1987. 48. A. K. Jonscher, Dielectric Relaxation in Solids, Chelsea Dielectrics, London, 1983. 49. K. Funke and C. Cramer, Conductivity spectroscopy, Curr. Opin. Solid State Mater. Sci, 2, 483–490 (1997). 50. J. Van Turnhout, Thermally stimulated discharge of electrets, in Electrets, G. M. Sessler (Ed.), Topics in Applied Physics Vol. 33, Springer, Berlin, 1980. 51. P. Pissis, A. Anagnostopoulou-Konsta, L. Apekis, D. Daoukaki-Diamanti and C. Christodoulides, Dielectric studies of water in water-containing systems, J. Non-Cryst. Solids, 131–133, 1174–1181 (1991). 52. C. Christodoulides, Determination of activation energies by using the widths of peaks of thermoluminescence and thermally stimulated depolarization currents, J. Phys. D, 18, 1501–1510 (1985). 53. A. Kyritsis, P. Pissis, J. L. Gomez Ribelles, and M. Monleon Pradas, Depolarization thermocurrent studies in poly(hydroxyethyl acrylate)/water hydrogels, J. Polym. Sci. Part B Polym. Phys., 32, 1001–1008 (1994). 54. J. M. Meseguer Duenas, D. Torres Escuriola, G. Gallego Ferrer, M. Monleon Pradas, J. L. Gomez Ribelles, P. Pissis, and A. Kyritsis, Miscibility of poly(butyl acrylate)-poly(butyl methacrylate)sequential interpenetrating polymer networks, Macromolecules, 34, 5525–5534 (2001). 55. P. Pissis, A. Kyritsis, J. M. Meseguer Duenas, M. Monleon Pradas, D. Torres Escuriola, G. Gallego Ferrer, and J. L. Gomez Ribelles, Dielectric and dynamic mechanical studies on homogeneous PBA/PBMA interpenetrating polymer networks, Macromol. Symp., 171, 151–162 (2001). 56. A. Espadero Berzosa, J. L. Gomez Ribelles, S. Kripotou, and P. Pissis, Relaxation spectrum of polymer networks formed from butyl acrylate and methyl methacrylate monomeric units, Macromolecules, 37, 6472–6479 (2004). 57. A. Kyritsis, J. L. Gomez Ribelles, J. M. Meseguer Duenas, N. Soler Campillo, G. Gallego Ferrer, and M. Monleon Pradas, α-β splitting region in the dielectric relaxation spectrum of PEA-PEMA sequential IPNs, Macromolecules, 37, 446–452 (2004).

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58. J. A. Gomez Tejedor, T. Rodriguez Acosta, J. L. Gomez Ribelles, G. Polizos, and P. Pissis, Poly(ethyl methacrylateco-hydroxyethyl acrylate) random co-polymers: Dielectric and dynamic-mechanical characterization, J. Non-Cryst. Solids, 353, 276–285 (2007). 59. G. Gallego Ferrer, M. Monleon Pradas, J. L. Gomez Ribelles, and P. Pissis, Swelling and thermally stimulated depolarization currents in hydrogels formed by interpenetrating polymer networks, J. Non-Cryst. Solids, 235–237, 692–696 (1998). 60. A. Valles Lluch, A. Campillo Fernandez, G. Gallego Ferrer, and M. Monleon Pradas, Bioactive scaffolds mimicking natural dentin structure, J. Biomed. Mater. Res. B, 90, 182–194 (2009). 61. G. Williams, Molecular aspects of multiple dielectric relaxation processes in solid polymers, Adv. Polym. Sci., 33, 59–92 (1979). 62. M. Beiner, Relaxation in poly(alkyl methacrylate)s: crossover region and nanophase separation, Macromol. Rapid Commun., 22, 869–895 (2001). 63. R. Bergman, F. Alvarez, A. Alegria, and J. Colmenero, The merging of the dielectric α- and β-relaxations in poly-(methyl methacrylate), J. Chem. Phys., 109, 7546–7555 (1998). 64. V. A. Bershtein, L. M. Egorova, P. N. Yakushev, P. Pissis, P. Sysel, and L. Brozova, Molecular dynamics in nanostructured polyimide-silica hybrid materials and their thermal stability, J. Polym. Sci. Part B Polym. Phys., 40, 1056–1069 (2002). 65. G. P. Johari and M. Goldstein, Viscous liquids and the glass transition. II. Secondary relaxations in glasses of rigid molecules, J. Chem. Phys., 53, 2372–2388 (1970). 66. S. Kripotou, P. Pissis, E. Kontou, A. M. Fainleib, O. Grigoryeva, and I. Bey, Polycyanurate networks modified by polyoxytetramethylene glycol, Polym. Bull., 58, 93–104 (2007). 67. C. A. Angell, Relaxation in liquids, polymers and plastic crystals – strong/fragile patterns and problems, J. NonCryst. Solids, 131–133, 13–31 (1991). 68. A. Kyritsis, P. Pissis, and J. Grammatikakis, Dielectric relaxation spectroscopy in poly(hydroxyethyl acrylates)/water hydrogels, J. Polym. Sci. Part B Polym. Phys., 33, 1737–1750 (1995). 69. D. W. Schaefer and R. S. Justice, How nano are nanocomposites?, Macromolecules, 40, 8501–8517 (2007). 70. C. Sanchez, G. J. de A. A. Soler-Illia, F. Ribot, T. Lalot, C. R. Mayer, V. Cabuil, Designed hybrid organic-inorganic nanocomposites from functional nanobuilding blocks, Chem. Mater., 13, 3061–3083 (2001). 71. A. C. Balazs, Predicting the morphology of nanostructured composites, Curr. Opin. Solid State Mater. Sci, 7, 27–33 (2003). 72. D. Brown, V. Marcadon, P. Mele, and N. D. Alberola, Effect of filler particle size on the properties of model nanocomposites, Macromolecules, 41, 1499–1511 (2008). 73. V. M. Litvinov, H. Barthel, and J. Weis, Structure of a PDMS layer grafted onto a silica surface studied by means of DSC and solid-state NMR, Macromolecules, 35, 4356–4364 (2002). 74. C. S. Chouchaoui and M. L. Benzeggagh, The effect of interphase on the elastic behavior of glass/epoxy bundle, Compos. Sci. Technol., 57, 617–622 (1997). 75. F. He, L.-M. Wang, and R. Richert, Dynamics of supercooled liquids in the vicinity of soft and hard interfaces, Phys. Rev. B, 71, 144205 (2005). 76. F. W. Starr, T. B. Schroder, and S. C. Glotzer, Molecular dynamics simulation of a polymer melt with a nanoscopic particle, Macromolecules, 35, 4481–4492 (2002). 77. P. Scheidler, W. Kob, and K. Binder, The relaxation dynamics of a supercooled liquid confined by rough walls, J. Phys. Chem. B, 108, 6673–6686 (2004). 78. M. Alcoutlabi and G. B. McKenna, Effects of confinement on material behaviour at the nanometer size scale, J. Phys.: Condens. Matter, 17, R461–R524 (2005). 79. O. Becker and G. P. Simon, Epoxy layered silicate nanocomposites, Adv. Polym. Sci., 179, 29–82 (2005). 80. S. Merabia, P. Sotta, and D. Long, Heterogeneous nature of the dynamics and glass transition in thin polymer films, Eur. Phys. J. E, 15, 189–210 (2004). 81. G. Tsagaropoulos and A. Eisenberg, Dynamic mechanical study of the factors affecting the two glass transition behavior of filled polymers. Similarities and differences with random ionomers, Macromolecules, 28, 6067–6077 (1995).

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82. K. U. Kirst, F. Kremer, and V. M. Litvinov, Broad-band dielectric spectroscopy on the molecular dynamics of bulk and adsorbed poly(dimethylsiloxane), Macromolecules, 26, 975–980 (1993). 83. V. M. Litvinov and H. Spiess, H-2 NMR study of molecular motions in polydimethylsiloxane and its mixtures with aerosil, Makromol. Chem. – Makromol. Chem. Phys., 192, 3005–3019 (1991). 84. V. Arrighi, I. McEwen, H. Qian, and M. Prieto, The glass transition and interfacial layer in styrene-butadiene rubber containing silica nanofiller, Polymer, 44, 6259–6266 (2003). 85. L. Matejka, O. Dukh, and J. Kolarik, Reinforcement of crosslinked rubbery epoxies by in-situ formed silica, Polymer, 41, 1449–1459 (2000). 86. V. Arrighi, J. Higgins, A. Burgess, and G. Floudas, Local dynamics of poly(dimethyl siloxane) in the presence of reinforcing filler particles, Polymer, 39, 6369–6376 (1998). 87. J. Berriot, H. Montes, F. Lequeux, D. Long, and P. Sotta, Evidence for the shift of the glass transition near the particles in silica-filled elastomers, Macromolecules, 35, 9756–9762 (2002). 88. D. Fragiadakis, P. Pissis, and L. Bokobza, Glass transition and molecular dynamics in poly(dimethylsiloxane)/silica nanocomposites, Polymer, 46, 6001–6008 (2005). 89. D. Fragiadakis, P. Pissis, and L. Bokobza, Modified chain dynamics in poly(dimethylsiloxane)/silica nanocomposites, J. Non-Cryst. Solids, 352, 4969–4972 (2006). 90. D. Fragiadakis and P. Pissis, Glass transition and segmental dynamics in poly(dimethylsiloxane)/silica nanocomposites studied by various techniques, J. Non-Cryst. Solids, 353, 4344–4352 (2007). 91. E. Laredo, A. Bello, and M. Grimau, The analysis of TSDC peaks with a KWW relaxation function or a distribution of relaxation times in polymers, Polym. Bull., 42, 117–124 (1999). 92. P. Pissis, Molecular dynamics in thermoset nanocomposites, in Thermoset Nanocomposites for Engineering Applications, R. Kotsilkova (Ed.), Rapra, Shawbury (UK), 2007. 93. J. Mijovic, H. Lee, J. Kenny, and J. Mays, Dynamics in polymer-silicate nanocomposites as studied by dielectric relaxation spectroscopy and dynamic mechanical spectroscopy, Macromolecules, 39, 2172–2182 (2006). 94. K. A. Page and K. Adachi, Dielectric relaxation in montmorillonite/polymer nanocomposites, Polymer, 47, 6406–6413 (2006). 95. M. Moniruzzaman and K. I. Winey, Polymer nanocomposites containing carbon nanotubes, Macromolecules, 39, 5194–5205 (2006). 96. D. Stauffer and A. Aharony, Introduction to Percolation, Taylor and Francis, London, 1994. 97. D. S. McLachlan, C. Chiteme, C. Park, K. E. Wise, S. E. lowther, P. T. Lillehel, E. J. Siochi, and J. S. Harrison, AC and DC percolative conductivity of single wall carbon nanotube polymer composites, J. Polym. Sci. Part B Polym. Phys., 43, 3273–3287 (2005). 98. B. E. Kilbride, J. N. Coleman, J. Fraysse, P. Foumet, M. Cadek, A. Drury, S. Hutzler, S. Roth, and W. J. Blau, Experimental observation of scaling laws for alternating current and direct current conductivity in polymer-carbon nanotube composite thin films, J. Appl. Phys., 92, 4024–4030 (2002). 99. S. Pegel, P. Poetschke, G. Petzold, I. Alig, S. M. Dudkin, and D. Lellinger, Dispersion, agglomeration, and network formation of multi-walled carbon nanotubes in polycarbonate melts, Polymer, 49, 974–984 (2008). 100. R. Kotsilkova, D. Fragiadakis, and P. Pissis, Reinforcement effect of carbon nanofillers in an epoxy resin system: rheology, molecular dynamics, and mechanical studies, J. Polym. Sci. Part B Polym. Phys., 43, 522–533 (2005). 101. Ch. Pandis, E. Logakis, M. Chorianopoulos, A. Spanoudaki, A. Kyritsis, V. Peoglos, P. Pissis, M. Micusik, I. Krupa, M. Omastova, J. Piontek, and P. Poetshke, Thermal and electrical characterization of polypropylene/carbon nanotube nanocomposites, in NSTI-Nanotech 2007, Vol. 2, 166–169 (2007). 102. L. Wang and Z.-M. Dang, Carbon nanotube composites with high dielectric constant at low percolation threshold, Appl. Phys. Lett., 87, 042903 (2005). 103. V. Peoglos, E. Logakis, Ch. Pandis, P. Pissis, J. Piontek, P. Poetschke, M. Micusik, and M. Omastova, Thermal and electrical characterization of multi-walled carbon nanotubes reinforced polyamide 6 nanocomposites, J. Nanostruct. Polym. Nanocomp., 3/4, 116–124 (2007.) 104. D. Chatterjee, K. Yurekli, V. G. Hadjiev, and R. Krishnamoorti, Single-walled carbon nanotube dispersions in poly(ethylene oxide), Adv. Funct. Mater. 15, 1832–1838 (2005).

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13 Solid-State NMR Spectroscopy of Multiphase Polymer Systems Antonio Mart´ınez-Richa Departamento de Quimica, Universidad de Guanajuato, Guanajuato, Mexico

Regan L. Silvestri Department of Chemistry, Baldwin-Wallace College, Ohio, USA

Solid-state nuclear magnetic resonance (NMR) has emerged as a relatively facile technique for the characterization of multicomponent polymer systems. In particular, it has proved to be an excellent technique for probing the molecular structure, conformation and dynamics of polymer chains at the interfaces between phases in various types of multicomponent polymer systems. The advantage of solid-state NMR is its ability to non-destructively probe not only the bulk of the polymer, but moreover its ability to selectively probe the interface or interphase. As such, the technique has been extensively exploited in the study of multicomponent polymer systems. The dynamics of adsorbed molecules on a mineral filler surface and their orientation inside the galleries can be studied. To achieve 13 C spectral resolution in the solid-state magic angle spinning (MAS), dipolar decoupling and cross-polarization are applied, which enables the study of individual carbon atoms directly with excellent resolution and sensitivity. The literature on cross-polarization in solid state NMR is growing steadily. A literature search for articles containing this concept in the last eight years produces 483 references using SciFinder, and 632 using Scopus. Review articles on the application of solid-state NMR to the structure and dynamics of crystalline and noncrystalline phases in polymers [1] and in polymer blends [2] appeared in 1998. This chapter reviews the principles of NMR involved in the cross-polarization experiment and in relaxation measurements, and illustrates the recent applications of this technique to multiphase polymer systems with selected examples that have appeared in the literature in the last eight years. Specific examples of applications of NMR techniques to amorphous (noncrystalline) and crystalline domains are described. Some spectroscopic findings have

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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been related to structural features and other properties such as mechanical performance, photovoltaic energy conversion and gas permeation.

13.1 Introduction to NMR Atomic nuclei have an inherent characteristic referred to as nuclear spin. The nuclear spin of a nucleus is characterized by an angular momentum, and the angular momentum of a nuclear spin is quantized. Accordingly, values for the nuclear spin quantum number I are quantized [3]. Nuclear Magnetic Resonance (NMR) exploits a strong external magnetic field to partially polarize the nuclear spins. Nuclei with a spin quantum number of I = 1/2 yield high resolution NMR spectra and are therefore particularly informative, as is evident in the study of multiphase polymer systems. Nuclei with a spin quantum number of I = 1/2 such as 1 H, 13 C, 19 F, 29 Si, 15 N and 31 P display spectra with unique peaks for each magnetically in-equivalent nuclei in the chemical structure [4]. As such, localized information can be gained about individual atomic sites in the chemical structure. A positively charged nucleus naturally aligns with an external magnetic field, and spins on its axis in parallel alignment with the external magnetic field. Tipped off its axis by a radio frequency (rf) pulse, the spinning nucleus precesses in the magnetic field back to equilibrium alignment with the external magnetic field as shown in Figure 13.1. As the local magnetic field surrounding the nucleus shields the nucleus from the external magnetic field, the rate of this precession is governed by the amount of the shielding effect. The rate of precession of the nucleus is measured as the chemical shift, and is displayed along the x-axis in the spectrum. Accordingly, magnetically in-equivalent atoms appear as unique peaks in the NMR spectrum, yielding localized information at individual atomic sites within the chemical structure. In solution, local magnetic fields experienced by any NMR active nucleus are averaged by rapid isotropic motions which are faster than the timescale of the NMR experiment. As a result, relatively sharp NMR peaks are observed for polymer samples in solution. However, in the interest of using multiphase polymer systems as engineering materials, there is a need for them to be studied as solids. Foremost of interest is the collection of NMR information in the solid state for the 13 C nucleus.

Processional orbit

Nuclear magnetic dipole μ

Spinning proton

B0

Figure 13.1 Precession of a nuclear spin tipped off its axis back to equilibrium alignment with the external magnetic field. Reprinted from [5]. Copyright (2005) with permission from John Wiley & Sons.

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Observation of the 13 C nucleus in the solid state, however, is complicated by line broadening caused by strong dipolar interactions between the nearly 100% abundant 1 H isotope and the 1.1% naturally abundant 13 C isotope. Heteronuclear dipolar decoupling reduces this line broadening by a high-powered radio frequency pulse at the 1 H frequency during the time which the 13 C signal is observed [6]. The collection of NMR spectroscopic data in the solid state is also complicated by the high chemical shift anisotropy which is not motionally averaged in the solid state to the extent that it is in the liquid state. A common technique termed magic angle spinning (MAS) is used to reduce the line broadening caused by chemical shift anisotropy in solids. Magic angle spinning consists of spinning the solid sample at a rate much ˚ higher than the 13 C chemical shift anisotropy, typically a few thousand hertz, at an angle of precisely 54.7A relative to the static magnetic field. At this so called magic angle, the broad chemical shift anisotropy pattern is reduced to a single peak at the isotropic chemical shift [7]. The final complication concerning the collection of 13 C NMR data in the solid state is sensitivity. As mentioned, the 13 C isotope is only 1.1% naturally abundant. With the above-mentioned heteronuclear dipolar coupling and chemical shift anisotropy consideration causing this inherently weak signal to be spread along a wide line-width that can only be partially reduced, the weak signal due to the low relative abundance of the 13 C isotope results in a low signal-to-noise ratio. A technique known as cross polarization (CP) is used to enhance the 13 C signal via transfer of polarization from the nearly 100% abundant 1 H spin reservoir. Cross polarization is accomplished via simultaneously applying spin-locking rf pulses to the 13 C and 1 H nuclei. Under special conditions when both nuclei precess at the same frequency, referred to as the Hartmann-Hahn match, magnetization is transferred from the abundant 1 H spin reservoir to the dilute 13 C spin reservoir. Experimentally, the two nuclei are spin-locked at the Hartmann-Hahn frequency match conditions for the appropriate contact time in the pulse sequence to allow magnetization to be transferred from the abundant to the dilute spin reservoir, thus resulting in signal enhancement of the more dilute 13 C spin reservoir which is observed [8]. The combination of these three techniques (dipolar decoupling, magic angle spinning and cross polarization) into one experiment provides a routine method for the collection of high-resolution 13 C NMR spectra of solids. Figure 13.2 shows progressively the effects of these techniques as they are applied in combinations and ultimately simultaneously in one experiment [9]. The technique combining dipolar decoupling, magic angle spinning and cross polarization is routinely applied to solid samples, and applications to polymer systems are sufficiently plentiful to fill entire professional reference textbooks [10, 11]. Regarding multiphase polymer systems, applications of the solid-state NMR technique are highly informative for (a) conformational elucidation of solid polymers via analysis of chemical shifts, and (b) the determination of crosslinking structures in solid polymers which are crosslinked. Modern NMR spectroscopy is carried out as a pulsed technique, where magnetization at equilibrium with the static magnetic field is perturbed from equilibrium via a radio frequency pulse. Sophisticated and elegant sequences of rf pulses are used to perturb the magnetization in specific ways. Observation of the magnetization as it precesses back to equilibrium via various clever sequences of rf pulses allows the collection of information far beyond the simple spectrum. The study of selectively pulsed magnetization as it relaxes back to equilibrium yields insight into both the structure and molecular dynamics of the system. Relaxation of a nucleus back to equilibrium is governed by fluctuations in the local magnetic field that the nucleus experiences. The local field around a nucleus is modulated from the static magnetic field due to the local environment around the nucleus. The local environment around the nucleus is dictated by the physical–chemical structure around the nucleus, and by molecular motions that the considered nucleus is involved in. Of particular interest herein is the study of relaxation as it gives information about phases and phase structure in multiphase polymer systems; relaxation is modulated by physical–chemical structure and molecular motions, both of which yield insight into phases and phase structure in multiphase polymer systems. Various relaxation parameters, including T1 , T2 and T1ρ

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(a)

(b)

(c)

C

(d)

C-O

O-CH3 CH2 -CH3

CH3 CH2 C C=O OCH3

Figure 13.2 The 13 C NMR spectra of poly(methyl methacrylate) in the solid state: (a) PMMA studied using experimental conditions for solution spectra; (b) dipolar decoupling with cross-polarization; (c) dipolar decoupling with magic angle spinning at 3 kHz; and (d) dipolar decoupling with magic angle spinning (3 kHz) and cross-polarization. Contact times in the range 1–10 ms were used. Reprinted from [9]. Copyright (1983) from CFC Press, Guelph, Ontario.

can be measured, which may yield information about molecular motions in different frequency ranges that occur over different scales of distance [12]. A traditional or 1-dimensional NMR spectrum is displayed as frequency in ppm along the x-axis and intensity of absorption along the y-axis. Two-dimensional NMR techniques add a second dimension to the NMR experiment, and as a result data are displayed with time or frequency along both the x and y axes and intensity along the z axis. Two-dimensional NMR techniques serve to enhance the resolution of a traditional 1-dimensional spectrum by spreading the peaks out in a second dimension, or to observe through bond or through space interactions, the latter being related to physical structure such as conformation [13]. While two-dimensional NMR techniques have become routine in solutions, they are still somewhat experimentally cumbersome to conduct in the solid state.

13.2 Phases in Polymers: Polymer Conformation Polymers which are semi-crystalline contain, in the solid state, amorphous (noncrystalline) and crystalline domains. These domains constitute separate phases that can be selectively studied by a variety of techniques

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CH2

C C φ φ 1 2

H

Figure 13.3

523

CH2 C

H

Definition of torsion angles φ i for a poly(ethylene) chain.

including solid-state NMR. The characteristics of the phases that exist in a semi-crystalline polymer material are a result of the internal organization of the individual chain molecules. In the amorphous phase, randomly entangled polymer chains are present. In the crystalline domains, the molecules pack together in an ordered array forming regular or crystalline arrangements. Several levels of structural organization can be recognized in a polymer material, meaning that one or several phases can be present at a time. These observed forms are known as the morphology of polymeric materials. The observance of various polymer phases is temperature-dependent. At higher temperatures, changes in structural organization and molecular motion take place, and new arrangements are present due to changes in thermodynamic properties of the system such as enthalpy and entropy. Also, crystallization conditions affect the crystalline/amorphous content of a polymer. Chain conformation along the polymer backbone is defined by the relative geometrical arrangement of the chemical groups and their orientations. Sizes and shapes of polymer chains (extended or folded geometries) are a consequence of particular conformational characteristics, which in turn are defined by torsional angles that change with backbone bond rotations. Many of the phenomena associated with changes in macromolecular physical properties are related to the conformational characteristics of the polymer chains. In that regard, absence or presence of one or more crystalline phases is linked to the existence of certain conformers along the polymer backbone, which at the same time are associated with torsion motions of many bonds. Conformation in general is described by torsion angles. Conformational analysis deals with the study of the different ways that molecules can change their geometries by torsion rotations about covalent bonds [14]. For example, a polymer chain composed of methylene carbons can adopt different conformations that depend upon the values of torsion angles φ i (see Figure 13.3). Due to the steric restrictions imposed upon torsion angles, potential energy curves in general have minima at φ i = 0, +120◦ and –120◦ , which are known as trans, gauche(+) and gauche(–) conformations, respectively. After considering the energy barriers between trans and gauche conformers, it can be computed that for poly(ethylene) each backbone bond is in average ≈60% trans and 40% gauche(+) and gauche(–), at room temperature. The population of gauche conformations increases at higher temperatures [15]. In the case of polymers with rings in the polymer backbone such as polysaccharides, the secondary structure is defined by the extent of rotation along the torsion angles ψ and  about the glycosidic linkages, and to a lesser extent by the covalent glycosidic linkage bond angle τ (see Figure 13.4) [11, 16]. As an illustration of the nature of conformation in these polymers, torsion angles for a (1→ 4)-α-D-glucan polysaccharide are shown in Figure 13.4. Barriers to bond rotations about the torsion angles ψ and  are high in polysaccharides [16]. Accordingly, polysaccharide conformation depends strongly on the environment surrounding the polymer chain, and a broad variety of molecular shapes have been found, for example, in cellulose derivatives [17]. The observance of extended single or multiple helical chains is related to changes in the secondary structure. The dependence of observed carbon-13 chemical shifts on conformation has been used to derive three-dimensional structures of other biological macromolecules, such as peptides and proteins [16]. High-resolution solid-state 13 C NMR provides important information on the polymer chain conformation, as observed chemical shifts and relaxation times are sensitive to chemical structure, morphology, orientation,

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Handbook of Multiphase Polymer Systems 6

CH2OH

O

O

φ

2

HO

τ

OH

O

OH

3

3´

HO

2´

Ψ

6´ CH

2OH

O

n

Figure 13.4 Chemical structure for a (1→ 4)-α-D-glucan polysaccharide. Torsion angles  and ψ and the covalent glycosidic linkage bond angle τ are defined as shown.

domain size, heterogeneity and the chain dynamics of materials. In contrast to solution NMR, in the solid state modulations to the local magnetic environment experienced by a nucleus due to surrounding nuclei are not averaged out by molecular motions. In most cases, differences in conformation between the amorphous and crystalline phases result in the observation of different peaks for the same nuclei (carbon-13 and other spin 1/2 nuclei) in the molecular structure. However, NMR chemical shifts do not depend upon the overall chain dimensions, but upon the local conformation that surrounds the nuclei being studied. The interpretation of spectra showing two or more peaks for the same nuclei has to be conducted with care, as there is not a simple unambiguous relationship between chemical shift and conformation. In the amorphous phase, the existence of a variety of conformations that do not interconvert rapidly enough on the NMR timescale translates into the observance of broad lines in the NMR spectra for polymers below their glass transition temperature. Conformational disorder, the distribution of interchange distances, the distribution in valence angle values and other effects that are locked into place in the glassy state result in a continuous distribution of 13 C chemical shift values. These effects result in inherently broadened resonance lines for which NMR peak narrowing techniques are not sufficient to resolve the lines of different species that are present in the amorphous phase. Crystalline phases, on the other hand, are ordered and the number of conformations is thereby restricted. As a consequence, there is more room for localized molecular motion in amorphous domains. For molecular motions that involve conformational transitions, such conformational transitions and molecular motions lead to differences in chemical shifts and relaxation times that can be measured by NMR techniques. Due to the nature of the amorphous phase, there is more room for molecular motion. This means that differences in relaxation times can be used to identify signals originating from apparently different morphologies. If the Newman projections of the gauche and trans conformations are analyzed, it can be inferred that ˚ than in carbons in the γ position are closer in the gauche conformer (intramolecular distance of around 3A) ˚ the trans conformer (intermolecular distance of around 4A, see Figure 13.5) [15]. As previously discussed,

γ

C

Figure 13.5

H

H

H

H

H

H

H

H

C

γ

C

C

(a)

(b)

Newman projections of (a) trans (φ = 0◦ ) and (b) gauche (φ = 120◦ ) conformers for poly(ethylene).

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many semi-crystalline polymers favor a larger population of more extended trans conformations to facilitate packing in crystalline arrangements. This generally means that the population of trans conformers is larger. The gauche arrangement results in a shielding effect for carbon-13 NMR chemical shifts of about –5 ppm relative to the chemical shift of the nuclei in the non-shielded trans conformer [15]. This variation in magnetic shielding between trans and gauche conformations, and the resulting variation in the observed chemical shift for nuclei in trans and gauche conformations, is referred to as the gamma-gauche effect. A classic example of the influence of the gauche effect on 13 C chemical shifts is observed in the signals for crystalline and amorphous phases of polyethylene. The planar trans zigzag conformation is confined in the tight packing of the crystalline phase in the glassy state, whereas in the amorphous phase rapid exchange between the gauche and trans conformations occurs. Consequently, the chemical shift observed for the amorphous phase is the average of both conformations, and is shifted upfield with respect to the chemical shift for the crystalline phase. In the case of nonolefinic polymers such as polysaccharides and other biological macromolecules, characterization of crystalline polymorphs and secondary structures thereof can be achieved using solid-state NMR as chemical shifts are very sensitive to molecular conformation [11, 15–17]. These observations, coupled with measurements of interatomic distances by rotational echo double resonance (REDOR), can be used to construct three-dimensional structures of macromolecules, a technique that has found particularly valuable applications for biological macromolecules [18].

13.3 High Resolution 13 C NMR Spectroscopy of Solid Polymers As many polymeric materials for practical applications are heterogeneous (multicomponent) systems, it is quite common to work with multiphase morphologies in the application of polymer science. As described previously, high-resolution solid state NMR spectra of rare spin systems such as carbon-13 and silicon-29 are usually recorded using cross-polarization (CP) between the observed nucleus and abundant proton spin, combined with magic angle spinning (MAS) and 1 H dipolar decoupling. Variable temperature CP-MAS experiments have been applied to study phase transitions, molecular motions and chemical exchange. In general, the intensity of crystalline peaks relative to amorphous peaks increases mainly due to (1) an increase in the amount of crystalline phase and (2) an increase in the cross-polarization efficiency for more crystalline structures. Other important consequence of phase separation is the effect on molecular dynamics. As discussed in the Section 13.1, the different domains present in polymer materials usually show differences in molecular mobility. These features can be detected by the measurement of relaxation parameters (such as T1 , T2 and T1ρ ). Nuclei in amorphous domains show a higher mobility than those measured for species located in the more rigid crystalline components (see also Section 13.6). The applications of high-resolution solid-state NMR spectroscopy to the study of solid polymers are numerous. To illustrate applications of this technique to the study of multiphase morphology, there follows selected examples from the previous eight years.

13.3.1

Chemical Shift

Chemical shift is a very important parameter in high-resolution NMR studies of polymers. Observed chemical shifts are related to the magnetic environment of the nucleus considered. They provide information not only on the chemical structure of the polymer, but also on conformation and higher-order arrangements such as crystal and noncrystalline (amorphous) structure. In that regard, for a single crystal the chemical shift depends

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on the angle between the crystal axis and the magnetic field, because the electronic distribution around the nucleus is not necessarily symmetric. As previously mentioned, amorphous polymers are composed of a broad range of conformers. One might anticipate that each of these species has a distinct chemical shift. However, due to the nature of magnetic heterogeneity which results in a continuous distribution of chemical shifts, in most cases only broad resonances are observed in the spectra. On the other hand, well-resolved peaks are generally observed for crystalline phases. Analysis of crystalline and noncrystalline peaks, along with spin-diffusion techniques, has been used to determine the degree of phase separation in semi-crystalline polymer systems. 13.3.2

Polyolefines

Polyethylene, the simplest possible polymer, is a typical example illustrating the phenomenon of phase separation that is due mainly to physical heterogeneity. The observed degree of crystallinity is a result of the ratio of trans and gauche conformations, and is influenced by many factors such as aging, processing conditions, method of crystallization, etc. Due to the gamma-gauche effect, carbons in an all-trans conformation resonate at a chemical shift of 33 ppm, whereas carbons in amorphous domains appear at around 30 ppm. From careful analysis of the line shapes and integration of these peaks, the ratio of crystalline to amorphous phase can be determined from the solid-state NMR spectrum. Quantitative measurements of crystallinity for ultradrawn polyethylene have been reported [19] and discussed by Mirau [11]. A modified method to determine the phase content in semi-crystalline polymers by solid-state NMR has also been recently proposed [20]. In this approach, a spin-diffusion (SD) filter is inserted in the pulse sequence prior to cross-polarization (CP) in a 13 C-CP-MAS experiment. This method was exploited to determine the degree of crystallinity of well-known polyolefine samples: one high-density polyethylene (HDPE) sample and four styrene-isoprene diblock copolymers, all of known phase contents. The calculated crystallinities are comparable with those obtained from wide-line 1 H NMR for contact times τ cp > 1 ms. Determinations using τ cp < 1 ms yield values which are approximately 10% smaller. The combined SD/CP-MAS technique is reliable under different Hartmann-Hahn conditions. It is expected that this technique can be applicable to any multiphase system in which spin-diffusion between different regions/phases exists and where the magnetization of different phases can be selectively filtered out. Analogously, solid-state NMR was used to study the molecular morphology of bioadhesive polymer blends based on crosslinked poly(acrylic acid) and starch. Comparison of CP-MAS spectra and corresponding molecular dynamics (see Section 13.5.4) between a sample obtained by a spray-dried formulation and that of a physical mixture reflects the differences in morphology [21]. 13.3.3

Polyesters

Carbon-13 CP-MAS NMR studies have revealed that crystallization conditions have only slight effects on the molecular structure and morphology of poly(L-lactide) PLLA, a biodegradable polyester [22]. Inclusion compounds of a poly(ε-caprolactone)/poly(L-lactide) diblock copolymer and α-cyclodextrin were studied by solid state NMR [23]. Carbon-13 solid-state NMR spectra of the diblock copolymer, αcyclodextrin, and the inclusion compound are shown in Figure 13.6. It can be seen that broader peaks are observed in the inclusion compound due to the different morphology. Based on the analysis of chemical shifts, it is concluded that the crystalline conformation of the PCL blocks in the inclusion compound is similar to that observed in PCL, and correspond to an all-trans planar zig-zag conformation. Phase separation was inferred from the analysis of values of the spin-lattice relaxation time in the rotating frame H T1ρ . Multiphase morphology has been reported for low-molecular weight PCLs for the first time. CP-MAS and MAS 13 C-NMR spectra from a single pulse excitation experiment for PCL obtained at 60◦ C are shown in

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(a)

ssb

ssb

(b)

(c)

240

220

200

180

160

140

120 δ

100

80

60

40

20 ppm

(13C)

Figure 13.6 13 C CP-MAS spectra of (a) bulk diblock PCL-b-PLLA copolymer; (b) pure α-CD; and (c) PCL-bPLLA-α-CD Inclusion Compound. Reprinted from [23]. Copyright (2005) with permission from John Wiley & Sons.

Figure 13.7. Peak patterns observed in the CP-MAS spectra for all PCLs are similar to those reported in the literature. However, MAS spectra show important differences, in particular for the signals due to the carbonyl and to the methylene carbon adjacent to oxygen. These differences indicate the coexistence of various types of phases, which are formed during crystallization. The occurrence of co-crystallization is rare in polymer science, and there are no previous reports of this phenomenon in PCLs [24]. Aliphatic–aromatic copolyesters of poly(butylene adipate-co-butylene terephthalate) [P(BA-co-BT)] with different copolymer compositions have been studied by solid-state NMR [25]. WAXD patterns indicate that copolyesters with 10–25 mol% BT contain PBA crystals, whereas those with 27.5–80 mol% BT contain PBT crystals. In contrast, solid state 13 C NMR analyses demonstrate that copolyesters with 20–30 mol% BT units contain both PBA and PBT crystals. The crystalline component in the lower amount cannot be detected by X-ray diffraction. Figure 13.8(a) and (b) show the high-resolution solid state Carbon-13 NMR spectra for P(BA-co10 mol%BT) in the regions around 130 and 34 ppm, respectively. The spectra (a) were obtained by

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Handbook of Multiphase Polymer Systems O

O g

O

HO

a

h

d i

f

OH

O

e

n

O

a

*

d,e

g f

h

i

*

180

160

140

120

100

80

60

40

20

ppm

Figure 13.7 Solid-state 13 C NMR (75.47 MHz) spectrum for poly(ε-CL) obtained with crude Y. lipolytica lipase in the presence of n-heptane in CD3 Cl. Reaction conditions: 3 mmol ε-CL/100 mg lipase. T = 60◦ C, t = 360 h reaction time, Mn(NMR) = 975 Da. CPMAS spectrum, contact time 1 ms, repetition time 3 s (bottom); MAS spectrum with a repetition time of 20 s (top). Spinning sidebands are indicated by an asterisk. Reprinted from [24]. Copyright (2008) with permission from John Wiley & Sons.

conventional CP/MAS 13 C NMR, and the less mobile component (crystalline component) is enhanced in spectra collected above the glass transition temperature of the polymer due to enhanced efficiency of transfer of cross-polarized magnetization for the crystalline component. The spectra (b) were obtained by the CPT1 pulse sequence, setting the 13 C spin-lattice relaxation decay time (t) to 3 s. It is expected that the less mobile crystalline component (which has a shorter 13 C T1 ) is much more enhanced in (b) as compared to the spectra

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Solid-State NMR Spectroscopy of Multiphase Polymer Systems (a) P(BA-co-10mol% BT) BT units PBT α-crystal

(b)

PBT amorphous

P(BA-co-10mol% BT) BA units PBA amorphous

PBA α-crystal

(a) CP/MAS

(if exist) crystalline enhanced

(a) CP/MAS

crystalline enhanced

(b) CPTI τ = 3s

(if exist) crystalline furthermore enhanced

(b) CPTI τ = 3s

crystalline furthermore enhanced

(c) 13C 45° single pulse pulse delay = 15s

amorphous enhanced

(c) 13C 45° single pulse pulse delay = 15s

140

135

amorphous enhanced

130 125 ppm from Me4Si

38

36

(c) P(BA-co-20mol% BT) BT units PBT α-crystal

PBT amorphous

P(BA-co-20mol% BT) BA units

PBA amorphous

crystalline enhanced

(a) CP/MAS

(b) CPTI τ = 3s

crystalline furthermore enhanced

(b) CPTI τ = 3s

amorphous enhanced

(b) 13C 45° single pulse pulse delay = 15s

135

32 30 ppm from Me4Si

PBA α-crystal crystalline enhanced

140

34

(d)

(a) CP/MAS

(c) 13C 45° single pulse pulse delay = 15s

529

130 125 ppm from Me4Si

38

crystalline furthermore enhanced

36

amorphous enhanced

34

32 30 ppm from Me4Si

Figure 13.8 Enlarged high-resolution solid-state 13 C NMR spectra for P(BA-co-10 mol%BT) and P(BA-co20 mol%BT) around 130 ppm ((a) and (c)) and 34 ppm (b and d), respectively. (a) CP-MAS; (b) using the Torchia sequence for determination of T1 , and setting the 13 C spin-lattice relaxation decay time (t) to 3 s; (c) 13 C 45◦ single pulse sequence with a pulse decay of 15 s. Reprinted from [25]. Copyright (2004) with permission from Elsevier.

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(a), as barring other considerations less mobile crystalline components generally have shorter 13 C T1 values. The spectra (c) were obtained by a 13 C 45◦ single pulse sequence (as opposed to a 90◦ single pulse sequence) with the recycle delay time set to 15 s, and mainly reflect the mobile (amorphous) component. In Figure 13.8, the 13 C chemical shifts for amorphous PBT, crystalline (α-crystal) PBT, amorphous PBA, and crystalline (α-crystal) PBA have been indicated by the dashed lines. The spectra (a)–(c) in Figure 13.8(b) clearly reveal that the BA units of P(BA-co-10 mol%BT) are located in both the crystalline and amorphous regions. On the other hand, as shown in the spectra (a)–(c) in Figure 13.8(a), only the resonance line assigned to the amorphous PBT is observed. Blends of poly (ethylene terephtalate) PET and natural rubber NR processed in a twin-extruder were analyzed by solid state NMR and SEM. Solid state CP-MAS 13 C NMR studies revealed an increase in the cross-polarization time of the carbonyl carbon and a decrease of the H T1ρ relaxation of the carbonyl groups in the PET/NR blend. These changes are associated with the interaction between the two phases, which leads to a toughening effect of NR on PET (see also Section 13.6) [26]. Polyhydroxyalkanoates (PHAs), bacterially-synthesized aliphatic polyesters, are naturally-occurring biodegradable and biocompatible polymers with many uses in the biomedical area. The composition and morphology of these materials and their blends were characterized by solid-state carbon-13 NMR spectroscopy some time ago [27]. In a more recent study, solid-state NMR techniques were employed to investigate structural characteristics of the semi-crystalline biopolymers poly(3-hydroxybutyrate) PHB and poly(3hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) containing 2.7 mol% and 6.5 mol% of hydroxyvalerate in different domains of the copolymer. Both the homopolymer PHB and its copolymers PHBV containing 2.7 mol% and 6.5 mol% HV are composed of amorphous and crystalline regions having distinct dynamics at temperatures above the glass transition temperature but below the melt temperature. The degree of crystallinity (χ c ) has been calculated by deconvolution of crystalline and amorphous components for the peak attributed to the methyl groups at 21.3 ppm using a pulse sequence without cross-polarization. The ratios of the crystalline domains or χ c values are about 68 ± 2%, 60 ± 2% and 56 ± 2%, for PHB, PHBV2.7 and PHBV6.5, respectively. Using different relaxation delays, separation of crystalline and amorphous components can be achieved (see Figure 13.9) [28].

13.3.4

Carbohydrates

The nature of the morphology of cotton fibers grafted with (a) poly(ethyl acrylate) PEA (g-PEA) and (b) poly(ethyl acrylate-block-styrene) (g-(PEA-b-PSty)) has been studied by CP-MAS solid state NMR [29]. In Figure 13.10, spectra for (a) cellulose, (b) cellulose grafted with PEA and (c) with PEA-b-PSty are depicted. The relative ratio of the crystalline to amorphous peaks (at 89 and 84 ppm, respectively) in the cotton spectrum (a) indicates that, as expected, a highly crystalline cellulose arrangement is present in this sample. In the cellulose grafted with PEA and PEA-b-PSty spectra, no changes in shape and peak position can be detected, which may indicate that the grafting process does not significantly change the original crystalline arrangement present in cellulose. From DSC and NMR investigations, evidence of a heterogeneous morphology for the graft-(PEA-block-PSty) skin is obtained, and this can be described as an inner layer of g-PEA sandwiched between a semicrystalline cellulose and polystyrene outer layer. Such morphology leads to a mobility restriction for PEA chains.

13.3.5

Conducting Polymers

Since the revolutionary discovery of conductivity in polyacetylene by Alan J. Heeger, Alan G. MacDiarmid and Hideki Shirakawa, for which they were awarded the Nobel Prize in 2000 for the discovery and development of conductive polymers, the study of conducting polymers has increased rapidly. Polythiophene derivatives

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B4 21.3 B3

B1 V1

*

(d)

* (c)

*

*

*

*

V3

B2 V4 V2

CH3 V5

19.8 (d)

(c)

(b)

*

(b)

* (a) (b – a)

(a) 175

150

125

100 75 ppm

50

25

0 27.5

25

22.5

20

17.5

15

ppm

Figure 13.9 13 C single pulse excitation MAS (SPEMAS) NMR spectra of (a) PHB with recycle delay of 2 s, (b) PHB (c) PHBV2.7; (d) PHBV6.5 with recycle delay of 30 s, and the subtraction of (b – a). Deconvolution of side chain methyl group peak indicates the existence of two components: one broad peak at 21.3 ppm associated with amorphous regions and a peak at 19.8 due to crystalline domains. Reprinted from [28]. Copyright (2007) with permission from Elsevier.

have been used as polymer photovoltaic devices due to their high durability and long lifetimes [30–32]. The properties of these polymers rely mainly on the existence of hydrogen-bonded networks. By means of Carbon-13 solid-state MAS using a symmetry-based double-quantum (2Q) dipolar recoupling technique, ˚ exist in the network, which is expected for it has been determined that carbon–carbon distances of 3.85A hydrogen-bonded carboxylic-acid derivatives of polythiophenes [33]. Changes in polymer morphology due to thermal degradation are linked to the decrease of device performance. CP-MAS studies allow the determination of changes in morphology due to degradation. Carbon-13 CP-MAS spectra of the thermocleaved polymer (polyacid), the starting material (non-cleaved polyester), and 3-carboxythiophene are shown in Figure 13.11. In the polyacid spectrum, peaks due to hydrogen-bonded species (at 168.6 ppm) and non-hydrogen-bonded species (at 164.1 ppm) can be distinguished. The intensity of signals for non-hydrogen-bonded entities increases in the thermocleaved (degraded) polymer. A new class of crosslinked proton-conducting membranes (CPMs) with polyoxyalkylene moieties has been prepared from poly-(styrene-co-maleic anhydride) modified with 2-aminoethanesulfonic acid sodium salt (AESA-Na) and polyoxyalkylenediamines (PEGDAs). Though amorphous, differences in morphology have been detected via chemical shifts and line broadening for the methylene carbons in the polyether block. The presence of sulfonic acid sodium salt causes a broader distribution of the polyether-segment environments and also reduces segmental motion of the polymer chains [34]. 13.3.6

Polymer Blends

The Flory-Huggins theory predicts that most polymer blends are immiscible, due mainly to the high degree of polymerization, which results in a small value for the combinatorial entropy of mixing. A polymer blend

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Handbook of Multiphase Polymer Systems C2 3,5 H OH 6

4 O HO 3

5

H O

H

2

C4

C1

O

H

C6

αy am

OH 1 H

(a)

(b)

* *

*

*

* *

*

(c)

250

200

150

100

50

0

ppm

Figure 13.10 Carbon-13 CP-MAS spectra of (a) cellulose; (b) and (c) grafted polymers CotEA2 and CotEA2Sty, respectively. Asterisks denote spinning side bands. Reprinted from [29]. Copyright (2007) with permission from American Chemical Society.

is considered miscible when a single glass transition Tg is seen in the DSC thermogram. However, this definition of a miscible polymer blend does not exclude phase separation at the molecular level. Phase separation occurring at the nanometric scale can be determined using solid-state NMR techniques. For quite some time now, numerous applications of solid-state NMR to mixtures of two or more polymers (polymer blends) have been reported. The resolved peaks in the NMR spectra can be used to determine the heterogeneity of a blend. Blend miscibility can be studied at the microscopic level by CP-MAS and dynamic measurements [35]. These effects have been studied in starch/polycaprolactone blends [36]. The ratio of crystalline (rich in double-helix content) to amorphous phases for starch in the blend can be determined from integration of the signals corresponding to the glycosidic C4 carbon (at 81.4 ppm) in the CP-MAS spectra. Effects of the poly(ethylene oxide) (PEO) molecular weight on the miscibility of poly(methacrylic acid) (PMAA) and poly(ethylene oxide) (PEO) blends have been reported. One important result of these studies is that the PEO crystalline phase is not present in all of the blends studied. Phase separation was also studied by measurements of T1ρ (H), T1 (H) and T1 (C) relaxation times. T1 (H) and T1 (C) relaxation time values indicate that restriction of chain mobility of both PEO and PMAA occurs. Observed differences are due to the amount of hydrogen-bonding interactions in the different blends [37]. Using this approach, miscibility in polymer blends of poly(vinylphenol) (PVPh) and poly(vinylpyrrolidone) (PVP) has been studied by solid-state nuclear magnetic resonance (NMR) spectroscopy [37]. Carbon-13 CPMAS NMR spectra of blends with different compositions are shown in Figure 13.12. The chemical shift of the carbonyl carbon and the hydroxyl substituted carbon of the phenolic ring (C-6) varies due to a

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(a)

*

O

S

ssb

ssb

OH

S n

(b)

ssb

ssb

13COOH

S

(c) ssb

ssb

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*

O S

S n

240

200

160

120 ppm

80

40

Figure 13.11 Carbon-13 CP-MAS spectra of (a) thermocleaved polymer; (b) 3-carboxythiophene; and (c) noncleaved polyester [33]. ssb indicates spinning side bands. Reprinted from [33]. Copyright (2008) with permission from Elsevier.

hydrogen-bonding intermolecular association. Results of chemical shift analysis indicate that PVPh is completely miscible with PVP over the entire range of compositions due to this strong hydrogen bonding [38]. The miscibility of poly(vinyl isobutyl ether) (PVIBE) and poly(ε-L-lysine) (ε-PL) blends, which depends on many factors in the processing of the blends, can be related to the degree of crystallinity of each component [39]. Both components are semicrystalline polymers. Carbon-13 CP-MAS spectra for several blends of different compositions are shown in Figure 13.13. The most readily noticeable difference is observed in the peak shape of the CαH signal of the ε-PL at 57.8 ppm. Using variable contact time experiments, contributions from crystalline and noncrystalline domains can be separated. Peak patterns for the CαH peak and those for the CHOCH2 group of PVIBE are very sensitive to the contact time used. These differences are mainly due to the fact that the 1 H spin-lattice relaxation time in the rotating frame (H T1ρ ) is shorter for the crystalline phase than for the noncrystalline domains. The degree of crystallinity for the polymers in the blend can be computed after separation of the contributions from the crystalline and noncrystalline components. As such, it was found that crystallinity for PVIBE decreases to half of its initial value after blending. On the other hand, it was found that the degree of crystallinity of ε-PL remains in general the same in the blends. These observations suggest that the growth of the crystalline phase of PVIBE is largely affected by blending with ε-PL, whereas crystalline domains for ε-PL are not significantly influenced by blending with PVIBE. The only exception was observed for the PVIBE/ε-PL = 10/1 blend, where the excess of PVIBE impedes the crystallization of ε-PL. Crystal size domains for ε-PL in the blends are in the range of 50–100 nm.

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Figure 13.12 13 C CPMAS spectra at room temperature for PVPh/PVP blends: (a) 100/0; (b) 80/20; (c) 60/40; (d) 40/60; (e) 20/80; and (f) 0/100. Reprinted from [38]. Copyright (2001) with permission from American Chemical Society.

Polymers blends obtained from thermosetting resin and self-assembled amphiphilic block copolymers yield unique nanostructured morphologies (see also discussion in Section 13.5.1). An improved solid-state NMR method, along with other techniques, has been applied to characterize the phase behavior, miscibility, heterogeneous dynamics and microdomain structure in unsaturated polyester thermosets blended with PEO–PPO–PEO block copolymers. Inhibition of PEO crystalline domains in the blend has been observed by both DSC and 1 H MAS NMR. This fact is also reflected in the 13 C CP-MAS spectra, where peak shape and intensity for methylene carbon of PEO at 70–72 ppm give information regarding phase separation in the blend [40].

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(–NHCH2CH2CH2CH2CH(NH)CO–)n (–CH2CH(OCH2CH(CH3)2–)n CH3 side-chain CH main-chain CH2 CH-O-CH2

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Figure 13.13 Observed carbon-13 CP-MAS NMR spectra of pure PVIBE, pure ε-PL and the PVIBE/ε-PL blends at weight ratios of 10/1, 10/3 and 10/5. Expanded spectra of the aliphatic region from 45 to 95 ppm are shown as inlays. Reprinted from [39]. Copyright (2007) with permission from Atsushi Asano.

13.3.7

Interactions Between Polymers and Low-molecular Weight Compounds

Miscibility at the molecular level is important for some applications of polymer materials in the pharmaceutical area. As such, solid-state NMR has been applied to study the level of drug–polymer interactions in solid dispersion formulations. However, high quantities of drug are needed in order to study molecular interactions due to the insensitivity of the technique because the amount of material at the interface is inherently low. Using this approach, the nature of interactions for the ketoprofen–poly(ethylene oxide) PEO system has been studied. If a melt process is used to obtain a polymer–drug blend, the level of hydrogen bonding originally present in the crystalline ketoprofen is destroyed [41]. In contrast, using this melt process approach no bonding interactions were observed on both components of paclitaxel/poly (styrene-isobutylene-styrene) SIBS formulations. Even though the observed solid state NMR carbon-13 chemical shifts of paclitaxel do not change after mixing, some changes in molecular mobility (differences in carbon spin-lattice relaxation times) have been reported (see also Section 13.5.4) [42].

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Figure 13.14 Solid state 13 C NMR spectra of microbial poly(ε-L-lysine). (a) Normal spectrum; (b) crystalline component; and (c) amorphous component. Pulse sequences used are conventional CPMAS for (a) T1CP with a delay time of 7 s for (b), and saturation recovery with a delay time of 200 ms for (c). Reprinted from [44]. Copyright (2003) with permission from Elsevier.

13.3.8

Miscellaneous Polymers

Polypeptides are readily amenable to study by solid-state NMR techniques. They are model biomolecules for proteins as they can acquire the α-helix, β-sheet, ω-helix or other conformation under certain conditions. Differences in the range of 2–7 ppm (as a function of conformation) can be observed in 13 C chemical shifts [43]. As introduced earlier, poly(ε-L-lysine) is a water-soluble and biodegradable polymer that exhibits antibacterial activity [44]. The four methylene groups in the polymer backbone give the polymer a high degree of conformational flexibility. According to this, at least two different crystalline forms can exist. In the stable α-form, molecules have a trans-zigzag conformation and intermolecular hydrogen bonds are formed in an antiparallel sense, whereas the less stable γ -form involves intermolecular hydrogen bonds in a parallel direction. In Figure 13.14 the solid-state carbon-13 NMR spectra for the polymer (Figure 13.13(a)) and for the amorphous and crystalline components ((b) and (c) respectively, separated according to their relaxation times) are shown. Chemical shifts of the amorphous component are shielded, so they are observed upfield with respect to the crystalline peaks. The upfield shifts and the broadening of the methylene resonances are due to the conformational heterogeneity and γ -gauche effects in the amorphous component.

13.4 Additional Nuclei As stated previously, high-resolution NMR of some rare nuclei (such as carbon, nitrogen and silicon) is possible due to the simultaneous use of high power decoupling (to eliminate the broading resulting from the dipolar couplings with abundant nuclei such as 1 H) and magic angle spinning (MAS). However, highresolution solid-state NMR of polymers containing several abundant nuclei is damped due to the existence

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of complex pattern interferences provoked by their large magnetic dipole moments. Recently, some new techniques have been developed to overcome these difficulties. The Discrimination Induced by Variable Angle Minipulse (DIVAM) sequence [45–47] can separate signals on the basis of either relaxation or spin dynamics properties. It is composed of 12 low-amplitude pulses followed by a 90◦ observation pulse. For small minipulse angles, selection is achieved through relaxation differences, while at larger angles, selection is based on differences in the size of the chemical shielding anisotropy. Peaks for the different phases can be achieved using the appropriate sequence: narrow amorphous components nutate with the minipulse angle, while the crystalline signal lags behind and, sometimes, never truly nutates. Each component of the signal will therefore invert (cross-zero) at different minipulse angles. In this manner, one can determine the number of components that compose the signal and whether they exhibit amorphous or crystalline behavior. The morphological behavior of the poly[bis(trifluoroethoxy)phosphazene] PBFP, a highly hydrophobic and semicrystalline inorganic polymer with partially fluorinated side chains, has been studied using this approach [48]. This polymer contains 13 C, 1 H and 19 F nuclei in side chains, and 31 P in the polymer backbone. Using modified DIVAM nutation experiments and cross-polarization (CP) methods, identification of the different morphological components present in the 1 H, 19 F and 13 C signals was achieved (see Figure 13.15). In the 19 F spectrum of this polymer, only one signal at –77.6 ppm is observed as shown in Figures 13.15(b) and (c). However, peak deconvolution analysis indicates the presence of three components (one broad and two narrow lines). The narrow intense components correspond to the highly-mobile disordered amorphous phase of the polymer, and one broad component corresponds to its rigid crystalline phases. These assignments were made on the basis of the behavior of the linewidth with temperature where the amorphous signal is seen to broaden dramatically upon cooling. It was also determined that after heat treatment of solvent-cast PBFP, the degree of crystallinity is grossly increased from 70.0% to 80.0%. Similar results were obtained from the analysis of the 13 C, 1 H and 31 P spectra [48]. Morphology of materials that contain silica or organosilylated compounds can be characterized using solid-state silicon-29 NMR. The final morphology of hybrids obtained from tetraethoxysilane (TEOS) and 4-[(5-dichloromethylsilyl)pentyloxy] cyanobenzene (DCN) depends on the process used to obtain the hybrid. Different amounts of Q4 units (silicon atoms surrounded by a tetrahedron of OSi groups, Si(OSi)4 ) and D2 units (Si(R)(R )(OSi)2 , where R and R are alkyl groups) are obtained as a function of the TEOS/H2 O/HNO3 molar ratio [49]. The morphology of silica particles in polydimethylsiloxane networks can also be studied using this technique [50]. Amounts of the Q4 , Q3 (Si(OSi)3 OH) and Q2 (Si(OSi)2 (OH)2 ) species are related to the form in which the silica particles were generated. Also, the transverse nuclear magnetic relaxation time, T2 , is affected by the nature of the physical and chemical crosslinks present in the filled polydimethylsiloxane. Solid-state silicon-29 NMR has also been employed to study the nature of silicon bonds in organic–inorganic nanocomposites based on octa(ethylcyclohexylepoxidedimethylsiloxy) silsesquioxane (a POSS epoxide) and a polyamide. The molecular structure of the polyhedral oligomeric silsesquioxane (POSS) and the properties of the nanocomposite interphase varied systematically as a function of the initial formulation [51]. Polymer–silica nanocomposites have been prepared using two approaches: (a) by free-radical polymerization of 2-hydroxyethyl methacrylate (HEMA) in the presence of HEMA-functionalized SiO2 nanoparticles (Type 1) and (b) by simultaneous in situ growth of the silica phase through the acid-catalyzed sol–gel polymerization of tetraethoxysilane (TEOS) (Type 2) [52]. Type 1 systems exhibit a particle-matrix morphology in which SiO2 particles tend to form aggregates. On the other hand, type 2 systems possess a finer morphology characterized by a very open massfractal silicate structure. This is also reflected by the presence of different silicon chemical environments between the two samples. The relative proportions of Q2 , Q3 , and Q4 species vary between the two samples. The type 1 system displays a much more condensed silicate structure (around 89%) with higher amounts of Q4 species than that recorded for type 2 (around 78%), for which the Q3 environment predominates.

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Figure 13.15 (a) Room temperature 19 F direct DIVAM nutation experiment of solvent-cast PBFP spinning at 25 kHz. (b) Room temperature deconvoluted 19 F solid-state NMR spectrum of solvent-cast PBFP spinning at 25 kHz. (c) Room temperature deconvolved 19 F solid-state NMR spectrum of heat-treated PBFP spinning at 25 kHz. (d) The 19 F direct DIVAM spectra for minipulse angles near the first zerocrossing point at 9, 10.8 and 12.6◦ respectively (top to bottom) [48]. Reprinted from [48]. Copyright (2005) with permission from Springer Science + Business Media.

13.5 NMR Relaxation The NMR spectroscopic properties of a nucleus are largely dictated by the particular local magnetic field around the nucleus. A nucleus is subject to the magnetic field of neighboring nuclei through dipolar interactions. The extent to which the magnetic properties of a nucleus are modulated by the magnetic field of neighboring nuclei depends on the intermolecular separation. The intermolecular distance, and in fact the

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T2 T1ρ T1 0

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Figure 13.16 The solid-state 13 C NMR relaxation parameters T1 , T2 and T1ρ (in seconds) and the corresponding correlation times (in seconds) for the molecular motions which they are sensitive to.

very make up of the local environment itself, changes as molecular motions occur. Molecular motions induce variations in the orientation of the internuclei vectors with respect to the static NMR magnetic field. The relaxation of magnetization back to equilibrium is facilitated by molecular motions, and therefore NMR relaxation techniques are valuable in the study of both structure and molecular dynamics. Information about interactions between nuclear spins that change as the environment changes due to motions can be obtained from studies of NMR relaxation parameters. Various relaxation parameters can be measured, which are related to molecular motions on varied time scales. The correlation time of a molecular motion is related to the distance or scale over which a molecular motion occurs. The NMR relaxation parameters that are most commonly measured are T1 , T2 and T1ρ . T1 is the so called spin-lattice relaxation time, T2 is the spin-spin relaxation time, and T1ρ is the spin-lattice relaxation time in the rotating frame. Relaxation parameters are routinely measured for both 13 C and 1 H, however the measurement of relaxation parameters for 13 C nuclei allows the selective observation of relaxation parameters for each carbon atom in the polymer structure individually, resulting in localized site-selective information about molecular motions at various carbons through the polymer repeat unit. Each of the relaxation parameters T1 , T2 and T1ρ is characteristic of molecular motions involving different frequency ranges and, accordingly, different length scales as shown in Figure 13.16. Therefore, spin-lattice relaxation characterized by the relaxation parameter T1 probes fast motions with frequencies in the MHz regime. Such fast motions are small-scale short-range motions which are typically internal to a molecule. Similarly, spin-lattice relaxation in the rotating frame characterized by the relaxation parameter T1ρ probes slower motions with frequencies in the kHz regime. Motions in the kHz regime occur over a larger distance range, and therefore correspond to motions such as long chain motions. Recall that it is change in dipolar interactions that modulates the local field when motion occurs. The theoretical relationships between the 13 C NMR relaxation parameters T1 , T2 and T1ρ in the solid state and correlation time τ c are shown in Figure 13.17 [53]. Relaxation parameters are measured in both solution and the solid state. One example of a relaxation study of a multiphase polymer system in solution is that of complexes of poly(methacrylic acid) and poly(ethylene oxide) [37]. Measurement of 1 H transverse relaxation times and 13 C spin-lattice relaxation times yielded information about molecular mobility which resulted in conclusions about molecular structure. The relaxation measurements demonstrated greatly restricted chain mobility for both poly(methacrylic acid) and poly(ethylene oxide) chains, which was attributed to hydrogen-bonding interactions creating an interpolymer complex in solution. However once again, in the interest of using multiphase polymer systems as engineering materials, there is a need for them to be studied as solids and therefore the collection of NMR relaxation information in the solid state is most enlightening. 13.5.1

NMR Relaxation in the Study of Polymer Blends

The blending of biodegradable polymers with thermoplastics has recently gained wide interest as a way to engineer the properties of thermoplastics. One such biodegradable material is poly(ε-L-lysine). The

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Figure 13.17 Dependence of the relaxation times T1 , T2 and T1ρ on the correlation time τ c of the molecular motions responsible for the relaxation, as predicted by molecular motions that result in changes in dipole–dipole interactions [53]. Region A is characteristic of a rigid lattice, region B of a nonrigid solid, region C of a viscous liquid, and region D of a nonviscous liquid. Reprinted from [53]. Copyright (1971) Academic Press, New York.

characterization of poly(ε-L-lysine) by solid-state 13 C NMR spectroscopy allows observation of the crystalline and amorphous regions separately. The degree of crystallinity for microbial poly(ε-L-lysine) was measured to be 63% from the T1 relaxation decay curves (data were fitted to a biexponential curve decay model M01 exp (–τ /T11 ) + M02 exp(–τ /T12 ), from which the percentage of crystallinity was derived). Amorphous carbons are characterized by T1 values of around 1 s, whereas carbons in the crystalline domains show T1 values in the range of 25–66 s [44]. Poly(ε-L-lysine) blended with poly(vinyl isobutyl ether) studied by 1 H spin-lattice relaxation allowed the observation of signals from the poly(vinyl isobutyl ether) and poly(ε-L-lysine) domains separately [39]. Domains of poly(ε-L-lysine) were estimated to be in the range of 50–100 nm, and additionally observation of 1 H spin-lattice relaxation curves resulted in conclusions regarding a relationship between crystallinity and miscibility in these semicrystalline polymer blends: the miscibility of poly(vinyl isobutyl ether)/poly(ε-Llysine) blends is affected by crystallinity, and the crystallinity of the poly(vinyl isobutyl ether) is affected by the blending of the poly(ε-L-lysine). A similar study of biodegradable blends of starch/polycaprolactone by 13 C T1 , 1 H T1 and 1 H T1ρ resulted in the observation of phase separation and determination of the length scale over which phase separation occurs. Phase separated versus homogeneous blends were obtained, depending on the mixing conditions [36]. NMR relaxation studies have also been used widely to investigate the influence of processing techniques and parameters on the resulting blends. Variations in blends of poly(ethylene glycol) and silica produced by the sol-gel process have been monitored as a function of poly(ethylene glycol) content [54]. 13 C T1 values

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were followed with change in poly(ethylene glycol) content, where longer T1 values are indicative of more rigid chains on the right side of the T1 minimum as pictured in Figure 13.17. When the poly(ethylene glycol) content was low, a single relatively long T1 value was observed indicating that there are some motional restrictions on the poly(ethylene glycol) chains [54]. As the amount of poly(ethylene glycol) increased, two 13 C T1 values were observed: a relatively long 13 C T1 value similar to the one observed at low poly(ethylene glycol) contents, plus an additional shorter 13 C T1 value which was ascribed to relaxation in a region of the blend rich in poly(ethylene glycol). Proton spin-lattice relaxation in the rotating frame was also monitored for this system as a function of the blend composition. Observation of a single 1 H T1ρ value is taken as indication of homogeneity in a multicomponent polymer system, whereas the observation of two discrete 1 H T1ρ values is attributed to relaxation signals from inhomogeneous separated phases. Such conclusions can be drawn from 1 H T1ρ values by recalling that spin-lattice relaxation in the rotating frame probes motions which occur over larger distances. Relaxation values describe the rate of decay of magnetization as a function of time, by exploiting the appropriate pulse sequence. The rate of decay is exponential with time provided that the polymer under study displays a single relaxation component, and the relaxation parameter T1 , T2 or T1ρ describing the decay rate is deduced from the slope of the straight line plot which is obtained when the magnetization is plotted on the y-axis on a natural log scale. Figure 13.18 shows plots of the natural log of magnetization in a 1 H T1ρ experiment for these poly(ethylene glycol)/silica blends with various compositions. Note how on a natural log scale a straight line decay is observed for one composition which is described by a single 1 H T1ρ value and indicates phase homogeneity, whereas for other compositions a dual exponential decay is observed which is described by two 1 H T1ρ values and indicates inhomogeneous separated phases. As such, the crystallization behavior and microstructure of such blends can be elucidated. With the relatively recent discovery of conductivity in organic conjugated polymers, these materials have grown to be of great interest, and are also amenable to study by NMR relaxation techniques. The nanomorphology of conjugated polymer blends has been studied by 1 H T1 and 1 H T2 relaxation techniques to

5 4

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Figure 13.18 Plots of natural log of magnetization in a 1 H T1ρ experiment for sol-gel processed polymer blends of poly(ethylene glycol)-silica for samples with various PEG/TEOS compositions: () pure PEG, () sample 3, () sample 4, (•) sample 6, (◦) sample 7, and () sample 8. Five milliliters of TEOS and 0.1–0.8 grams of PEG were used. Reprinted from [54]. Copyright (1998) with permission from John Wiley & Sons.

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investigate the homogeneity of such materials, to estimate the size of phase separated domains and to compare local segmental mobility [55]. The 1 H T1 and 1 H T2 relaxation parameters were monitored with variations in blend composition and film processing techniques to evaluate the various materials produced. With the recent surge of interest in nanocomposite materials such as these and others, NMR spectroscopy has emerged as a perfectly suited technique. Due to the timescales of various NMR experiments, results concerning structure and morphology on a nanoscale can be obtained. Another recent example is a study of nanostructured segmented networks and blends of poly(dioxolane)/poly(methyl methacrylate) by 1 H T1 and 1 H T1ρ which revealed poly(dioxolane) domain sizes ranging between 1 and 20 nm [56]. Similarly, dimensions below 2–3 nm have been observed in the amorphous phase in blends of poly(vinylphenol) and poly(vinylpyrrolidone) prepared by solution-casting using 1 H T1ρ relaxation [38]. Furthermore, 13 C chemical shifts as described in Sections 13.3 and 13.4.1 were used to identify hydrogenbonding interactions. As stated previously, the advantage of 13 C NMR spectroscopy is that local information about individual carbon sites within the molecular structure is provided. Accordingly, the hydrogen-bonding interaction in this multiphase polymer system was identified to occur specifically between the hydroxyl group of poly(vinylphenol) and the carbonyl group of poly(vinylpyrrolidone). Referring back to chemical shift analysis which was described earlier, Figure 13.12 shows how in this example the chemical shift of the carbonyl carbon of the poly(vinylpyrrolidone) at 177 ppm increases with the increased content of poly(vinylphenol), and likewise the chemical shift of the hydroxyl-substituted carbon in the phenolic ring of poly(vinylphenol) at 153 ppm increases with the increased content of poly(vinylpyrrolidone). As another example, interactions between the two phases of poly(ethylene terephthalate) toughened by natural rubber via melt blending have been observed via 13 C chemical shift and 1 H T1ρ relaxation studies [26]. A decrease in the 1 H T1ρ values was interpreted as evidence of interactions between the carbonyl groups of poly(ethylene terephthalate) with some functionality such as hydroxyl groups in the natural rubber.

13.5.2

NMR Relaxation in the Study of Copolymers

Interesting and practically useful results concerning the NMR relaxation studies of multiphase polymer systems are also obtained in the field of copolymers. In a study of bacterial biopolymeric metabolites and the copolymer thereof, significant molecular mobility enhancement was observed for the copolymers relative to their respective homopolymers via monitoring 1 H T1ρ and 13 C T1 values [28]. Traditionally, butadienes are some of the most commercially interesting materials. Solid-state 1 H relaxation studies have been effective in evaluating the compatibilizing action of random copolymers of polystyrene and polybutadiene and poly(styrene-butadiene-styrene) triblock copolymers in incompatible polystyrene/polybutadiene blends [57]. 1 H T1 , 1 H T2 and 1 H T1ρ relaxation studies were performed on both compatibilized and uncompatibilized blends. Significant differences were found for blends compatibilized with triblock copolymers due to preferential localization of the copolymers at the polystyrene/polybutadiene interface. Segmented blocks in copolymers is a common strategy for engineering the mechanical properties of polymers by creating hard and soft segments. Solid-state NMR relaxation studies have been used to investigate the motional heterogeneity in poly(ether-block-amide) copolymers [58]. The 1 H T1ρ , 13 C T1ρ and 13 C T1 relaxation curves were monitored to study the motional heterogeneity in terms of content and length of hard/soft segments and microphase-separated morphology. In particular, the extent of the motional distribution of the two components was determined by relaxation measurements. A similar study was carried out on segmented polymer networks based on poly(N-isopropyl acryl amide) and poly(tetrahydrofuran) using 1 H T1 and 1 H T1ρ . The aim was to monitor the multiphase behavior of the segmented polymer networks as the nature of the polymerizable end group of the bis-macromonomer used for copolymerization was varied [59].

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Finally, copolymers grafted on the surface of fibers to be further used in a composite have also benefited from dynamic investigations by solid-state NMR relaxation studies [29]. Cotton fibers were modified by surface polymerization of ethyl acrylate followed by copolymerization with styrene. The graft copolymerencapsulated cotton fibers were studied by 1 H spin-lattice relaxation and the results showed a heterogeneous morphology of the grafted skin. Such fibers are intended to be used in polymer composite materials, which themselves benefit from study by NMR relaxation as described in the following section. 13.5.3

NMR Relaxation in the Study of Polymer Composites

Evidence gathered by NMR relaxation studies has been used to support a previously-proposed model for the structure of fibers of poly(p-phenylenebenzobisoxazole) [60]. Two component relaxation curves were observed and the short 1 H T1 and 1 H T1ρ values were attributed to the noncrystalline portion of poly(pphenylenebenzobisoxazole), whereas the long 1 H T1 and 1 H T1ρ values were attributed to the crystalline portion of the fiber. From this data the characteristic crystal size was determined to be of an order of 1 nm and a fiber crystal structure previously proposed by X-ray diffraction was confirmed. In a quite detailed study, the dispersion state of organomodified clay fillers in nanocomposites based on various thermoplastics was investigated by NMR relaxation [61]. The 1 H T1 saturation-recovery curves shown in Figure 13.19 for polyamide 6 and a nanocomposite of polyamide 6 with an organomodified clay demonstrate how the presence of clay shortens the relaxation time. The nanocomposites have a shorter relaxation time because of paramagnetically-induced relaxation at the polymer–clay interface, indicating nanodispersion of the clay. Through systematic study thereby, quantitative measurements of the degree of nanodispersion can be made. 13.5.4

NMR Relaxation in the Study of Polymers for Drug Delivery

Certainly one of the most exciting developments in polymer science recently has been the more advanced use of polymers through various strategies for the controlled delivery of drugs in pharmaceuticals. This area has also benefited from NMR relaxation studies. 1.0

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Figure 13.19 1 H T1 saturation-recovery curves for polyamide 6 and a nanocomposite of polyamide 6 with an organomodified clay showing how the presence of clay shortens the relaxation time. Reprinted from [61]. Copyright (2008) with permission from Elsevier.

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The drug paclitaxel is a coronary drug-eluting stent system incorporated in a nanocomposite polymer matrix [42]. 13 C T1 and 1 H T1ρ were used to demonstrate that the mobility of the paclitaxel is increased in the paclitaxel-poly(styrene-isobutylene-styrene) mixture. Moreover, a careful monitoring of the relaxation parameters can be used to understand that a change from crystalline packing to amorphous packing occurs in paclitaxel due to intermolecular interactions with poly(styrene-isobutylene-styrene). One of the most common methods for the exploitation of polymers in drug delivery is to use the polymer as a matrix for solid dispersion. This strategy creates drug–polymer solid solutions to enhance drug bioavailability for crystalline drugs with poor water solubility. Here, solid-state NMR relaxation measurements have been used to characterize the miscibility between poly(ethylene oxide) and ketoprofen, a model for crystalline drugs [41]. It was found that ketoprofen formed a complete molecular dispersion within the amorphous domains of poly(ethylene oxide). As a result, ketoprofen has a high molecular mobility in the blend which leads to an enhanced dissolution rate for the ketoprofen in aqueous media. Furthermore, hydrogen bonding was identified by solid-state NMR spectroscopy between the carboxylic group of ketoprofen and the ether oxygen of poly(ethylene oxide). The observation of this hydrogen bonding was used in explanation of the high miscibility that is observed between ketoprofen and poly(ethylene oxide). A more sophisticated use of polymers in drug delivery has been the exploitation of bioadhesive drug carriers. In one study, starch/poly(acrylic acid) blends were evaluated as bioadhesive drug carriers, and spraydried blends were compared with analogously equivalent simple physical mixtures [21]. Nanomorphology was investigated by 13 C NMR relaxation experiments and it was determined that spray-drying resulted in homogeneous blends, whereas phase separation into individual starch and poly(acrylic acid) molecular domains was observed for physical mixtures.

13.6 Spin Diffusion During an NMR experiment, magnetization may be transferred from one individual nucleus to another individual nucleus of the same type by spin diffusion. Dipolar interactions between neighboring nuclei induce spin flips that result in a transfer of magnetization between like nuclei yielding spin diffusion. In the case of 1 H nuclei which are highly coupled, and when T1  T2 , the dipolar couplings between 1 H nuclei are stronger than the coupling with the environment and the 1 H spin system reaches equilibrium by spin diffusion faster than it would otherwise reach equilibrium by relaxation [10]. Spin diffusion is described by the spin diffusion coefficient and a characteristic distance of transfer of the magnetization. NMR pulse sequences which are designed to measure the distance of magnetization transfer by spin diffusion are particularly useful for measuring spin diffusion distances between phases in multiphase polymer systems. A spin diffusion experiment is a general framework consisting of: 1. A first step, enabling to create a distribution of magnetization distinct from the one displayed at equilibrium. 2. A second step, letting 1 H spin diffusion take place over the sample during a given time tm . 3. Lastly, a third step, corresponding to the detection of the signal. During the first step, various experimental schemes may be used to induce a gradient of magnetization within the sample. One of the possible approaches is to vary the strength of the dipolar filter, as a function of the number of degrees cycles and delay time between consecutive 90 degrees pulses (this is the pulse sequence scheme of the so-called ‘dipolar filter’). In this way, information on domains related to strong 1 H–1 H dipolar couplings (‘rigid’ domains) can be filtered out.

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1.0

I/I0

0.8 0.6 0.4 0.2

stoichiometric 1H content

0.0 0

5

10

15

20

tm1/2[ms1/2]

Figure 13.20 Spin diffusion decay curves of PHEMA/PIB amphiphilic conetworks plotted as magnetization relative to the initial magnetization as a function of the square root of the mixing time. The plateau value following equilibration of the magnetization between the two phases due to spin diffusion can be used to calculate the stoichiometric proton ratio between the protons in the two phases. Different numbers of dipolar filters cycles (nc) were used: () nc = 1, (•) nc = 2, () nc = 3, () nc = 4, (䉬) nc = 6, (x) nc = 10 and (*) nc = 15. 10 μs spacing between the pulses were used. Reprinted from [62]. Copyright (2003) with permission from American Chemical Society.

Spin diffusion decay curves are plotted as magnetization relative to the initial magnetization against the square root of the mixing time. These plots contain three important parameters: (a) the plateau value; (b) the initial slope; and (c) the intercept of the extrapolated initial slope. Such spin-diffusion curves reach a plateau after a mixing time which is sufficiently long to allow the magnetization to equilibrate over the whole sample due to spin diffusion. The choice of the strength of the dipolar filter can be optimized via the plateau value. Once the magnetization is equilibrated between the two phases due to spin diffusion, the stoichiometric proton ratio between the protons in the two phases can be calculated from the plateau values. Average domain size of the mobile phase can be derived from the slope value. Using 1 H NMR spin diffusion experiments in the solid state, phase separation and morphology of conetworks have been investigated for a biocomponent nanophase-separated conetwork. Spin diffusion decay curves for PHEMA/PIB amphiphilic conetworks are shown in Figure 13.20 [62]. Nanophase-separated structures were found for both phases in the conetworks. The spin-diffusion experiments indicated a strongly separated morphology without any detectable interphase between the mixed components, as was deduced from the observation of an initial linear portion in the spin diffusion curves. Furthermore, the sizes of the hydrophilic and hydrophobic domains were determined to be in the 5–15 nm range magnetization relative to the initial magnetization as a function of the square root of the mixing time. Likewise, a similar 1 H NMR spin diffusion study of unsaturated polyester-based thermoset blends containing PEO-PPO-PEO block copolymers has been used to measure the interphase thickness in this multiphase polymer system [40]. Via 1 H NMR spin diffusion experiments in the solid state, conclusions were derived in this study concerning phase behavior, miscibility, heterogeneous dynamics and microdomain structure. In particular, by systematic comparisons between resin blends conclusions were made via this technique concerning the influence of binary polymer–polymer interactions on phase behavior, domain size and interphase thickness. Interphase thickness was found to be a key parameter in these systems. Interphase thickness was quantified by measuring the percentage of a component which is immobilized in the other component in the

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interphase region by spin diffusion. From this value, interphase thickness and domain size can be calculated. As such, the domain size was calculated to be 14 nm and the interphase thickness was calculated to be 0.3 nm. Via this same 1 H NMR spin diffusion experiment in the solid state which exploits a dipolar filter, information can be gained about crystalline and amorphous phases in multiphase polymer systems. An example is the characterization of ultradrawn polyethylene fibers by this technique to quantitatively determine crystallinity and domain sizes [20]. The 1 H spin diffusion experiments yielded amorphous domain sizes of 10 ± 5 nm and crystalline regions of 100 ± 50 nm diameters for these ultradrawn, ultrahigh molecular weight PE fibers. More interestingly, a second highly-mobile amorphous phase was detected and through investigation of the sample parameters was determined to be induced by the drawing process. Ultimately, five morphological components were identified in the ultradrawn, ultrahigh molecular weight PE fibers: crystal core (distinguishing orthorhombic and monoclinic), disordered all-trans interfacial and/or tie molecules, mobile amorphous regions, and highly mobile segments (probably at void surfaces). Moreover, the relative percentages of these components were quantified and the domain sizes were measured for each by the 1 H spin diffusion experiment. Following this work, further analogous and continuing work has recently been conducted on similar and additional multiphase polymer systems to quantify crystalline and amorphous phase content [20]. One example is the extension of this technique to the quantitative determination of crystallinity in styrene-isoprene diblock copolymers. The extension of this technique to describe multiphase morphology and determine domain size continues by adapting the technique through more sophisticated pulse sequences. For example, cross-polarization has been incorporated into the pulse sequence in certain applications, and also single-quantum and double-quantum dipolar filters have been exploited for selected applications [63].

13.7 Concluding Remarks The main experimental parameters connected to solid-state NMR spectroscopy which may be exploited in the study of multiphase polymer systems have been reviewed. In addition to the NMR techniques discussed herein, two-dimensional NMR spectroscopy can also yield valuable insights into multiphase polymer systems. Two-dimensional NMR spectroscopy, now used as a routine in the solution state, is quickly developing into a more routine technique in the solid state also. A comprehensive review of the current state of the art is given by deAzevedo et al., as may be applied to multiphase polymer systems [64]. The variety of NMR techniques available yields much information concerning the structure and dynamics of multiphase polymer systems. A review of the applications of NMR spectroscopy to the selective study of crystalline and noncrystalline phases is given by Yamanobe [1], and a review on the applications of NMR spectroscopy to the study of polymer blends and miscibility is given by Asano and Takegoshi [2]. In the study of multiphase polymer systems, solid-state NMR spectroscopy specifically yields valuable information concerning microstructural conformation of phases, miscibility and extent of phase separation at a molecular level, as well as phase domain size for phase separated systems.

References 1. T. Yamanobe, Structure and Dynamics of Crystalline and Non-Crystalline Phases in Polymers, in Solid-State NMR of Polymers, I. Ando, T. Asakura (Eds), Elsevier, Amsterdam, 267–306, 1998. 2. A. Asano and K. Takegoshi, Polymer Blends and Miscibility, in Solid-State NMR of Polymers, I. Ando, T. Asakura (Eds), Elsevier, Amsterdam, 351–414, 1998.

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3. A. Abragam, Principles of Nuclear Magnetism, Clarendon Press, Oxford, 1961. 4. J.C. Randall, in Polymer Characterization by ESR and NMR, A.E. Woodward and F.A. Bovey (Eds), ACS Symposium Series 142, American Chemical Society, Washington DC, page 93–118, 1980. 5. R.M. Silverstein, F.X. Webster and D.J. Kiemle, Spectrometric Identification of Organic Compounds, 7th Edition, John Wiley and Sons, Inc., 2005. 6. L.R. Sarles and R.M. Cotts, Phys. Rev., 111, 853 (1958). 7. E.R. Andrew, A. Bradbury and R.G. Eades, Nature, 182, 1659 (1958). 8. S.R. Hartmann and E.L. Hahn, Phys. Rev., 128, e 2042 (1962). 9. C.A. Fyfe, Solid State NMR for Chemists, CFC Press, Guelph, Ontario, 1983. 10. J.L. Koenig, Spectroscopy of Polymers, 2nd edition, Elsevier, New York, 1999. 11. P.A. Mirau, A Practical Guide to Understanding the NMR of Polymers, Wiley-Interscience, New York, 2005. 12. R.L. Silvestri and J.L. Koenig, Analytica Chimica Acta, 283, 997–1005 (1993). 13. R.R. Ernst, G. Bodenhausen and A. Wokaun, Principles of Nuclear Magnetic Resonance in One and Two Dimensions, Clarendon Press, Oxford, 1987. 14. H. R. Allcock, F. W. Lampe and J.E. Mark, Contemporary Polymer Chemistry, 3rd edition, Pearson Education, New Jersey, USA, 2003, Chapter 18: Conformational Analysis of Polymers. 15. A.E. Tonelli, Polymers From the Inside Out. An Introduction to Macromolecules, Wiley, New York, 2001. Chapter 5: The Conformational Characteristics of Polymers. 16. H. Saito, S. Tuzi and A. Naito, Polysaccharides and Biological Systems, in Solid-State NMR of Polymers, I. Ando, T. Asakura (Eds), Elsevier, Amsterdam, 891–921, 1998. 17. A.D. French and G.P. Johnson, Advanced Conformational Energy Surfaces for Cellobiose, Cellulose, 11, 449–462 (2004). 18. K. Werner, I. Lehner, H.K. Dhiman, C. Richter, C. Glaubitz, H. Schwalbe, J. Klein-Seetharaman and H.G. Khorana, Combined solid state and Solution NMR Studies of Alpha,Epsilon- 15 N Labeled Bovine Rhodopsin, J. Biomol NMR, 37, 303–12 (2007). 19. W-G. Hu and K. Schmidt-Rohr, Characterization of Ultradrawn Polyethylene Fibers by NMR: Crystallinity, Domain Sizes and a Highly Mobile Second Amorphous Phase, Polymer, 41 2979–2987 (2000). 20. L. Zhang, Z. Liu, Q. Chen, E.W. Hansen, Quantitative Determination of Phase Content in Multiphase Polymers by Combining Spin-Diffusion and CP-MAS NMR, Macromolecules, 40, 5411–5419 (2007). 21. D. Ameye, E. Pringels, P. Foreman, J.P. Remon, P. Adriaensens, L. Storme and J. Gelan, Correlation between the Molecular Morphology and the Biocompatibility of Bioadhesive Carriers Prepared from Spray-Dried Starch/Carbopolw Blends, Polymer, 46, 2338–2345 (2005). 22. A. Sroka-Bartnicka, S. Olejniczak, M. Sochacki, T. Biela and M.J. Potrzebowski, Solid State NMR Spectroscopy as a Tool Supporting Optimization of MALDI-TOF MS Analysis of Polylactides, J Am Soc Mass Spectrom, 20(1) 67–72 (2009). 23. F.E. Porbeni, I.D. Shin, X. Shuai, X. Wang, J.L. White, X. Jia and A.E. Tonelli, Morphology and Dynamics of the Poly(ε-caprolactone)-b-poly(L-lactide) Diblock Copolymer and its Inclusion Compound with α-Cyclodextrin: A solid-state 13 C NMR Study, J Polym Sci: Part B: Polym Phys, 43, 2086–2096 (2005). 24. K.A. Barrera-Rivera, A. Flores-Carre´on and A. Mart´ınez-Richa, Enzymatic Ring-Opening Polymerization of ε-caprolactone by a New Lipase from Yarrowia lipolytica, J. Appl. Polym. Sci., 109(2), 708–719 (2008). 25. Z. Gan, K. Kuwabara, M. Yamamoto, H. Abe and Y. Doi, Solid-State Structures and Thermal Properties of Aliphatic–Aromatic Poly(butylene adipate-co-butylene terephthalate) Copolyesters, Polymer Degradation and Stability, 83, 289–300 (2004). 26. P. Phinyocheep, J. Saelao and J.Y. Buzar´e, Mechanical Properties, Morphology and Molecular Characteristics of Poly(ethylene terephthalate) Toughened by Natural Rubber, Polymer, 48, 5702–5712 (2007). 27. Y. Inoue, Biodegradable Polymers, in Solid State NMR of Polymers, Ando I, Asakura T (Eds), Elsevier, Amsterdam, 771–817, 1998. 28. L. Zhang, H. Tang, G. Hou, Y. Shen and F. Deng, The Domain Structure and Mobility of Semi-Crystalline Poly(3-hydroxybutyrate) and Poly(3-hydroxybutyrate-co-3-hydroxyvalerate): A Solid-State NMR Study, Polymer, 48, 2928–2938 (2007).

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29. V. Castelvetro, M. Geppi, S. Giaiacopi and G. Mollica, Cotton Fibers Encapsulated with Homo- and Block Copolymers: Synthesis by the Atom Transfer Radical Polymerization Grafting-from Technique and Solid-State NMR Dynamic Investigations, Biomacromolecules, 8, 498–508 (2007). 30. M. Adi, T. Yohannes and T. Solomon, Solid-State Photoelectrochemical Device Based on poly(3-hexylthiophene) and an Ion Conducting Polymer Electrolyte, Amorphous Poly(ethylene oxide) Complexed with I3 − /I− Redox Couple, Sol. Energy Mater. Sol. Cells, 83(2/3) (special issue), 301–310 (2004). 31. G.G. Wallace, C.O. Too, D.L. Officer, and P.C. Dastoor, Photoelectrochemical Cells Based on Inherently Conducting Polymers, MRS Bull., 30(1) (special issue), 46–49 (2005). 32. E. Bundgaard and F.C. Krebs, Large-Area Photovoltaics Based on Low Band Gap Copolymers of Thiophene and Benzothiadiazole or Benzo-bis(thiadiazole), Sol. Energy Mater. Sol. Cells, 91(11) (special issue), 1019–1025 (2007). 33. M. Bjerring, J.S. Nielsen, A. Siu, N.C. Nielsen and F.C. Krebs, An Explanation for the High Stability of Polycarboxythiophenes in Photovoltaic Devices—A Solid-State NMR Dipolar Recoupling Study, Solar Energy Materials & Solar Cells, 92, 772–784 (2008). 34. P.L. Kuo, W.J. Liang, C.Y. Hsu and W.H. Jheng, Preparation, Characterization, and Properties of New Crosslinked Proton-Conducting Membranes with Polyoxyalkylene Moieties, Polymer, 49, 1792–1799 (2008). 35. H. Asano and K. Takegoshi, Polymer Blends and Miscibility, in Solid-State NMR of Polymers, I. Ando, T. Asakura (Eds), Elsevier, Amsterdam, 351–414, 1998. 36. J. Spevacek, J. Brus, T. Divers and Y. Grohens, Solid-state NMR Study of Biodegradable Starch/Polycaprolactone Blends, European Polymer Journal, 43, 1866–1875 (2007). 37. Z. Huipeng, W. Lin, G. Yang and Q. Chen, Molecular Weight Effect on the Complexation of Poly(methacrylic acid) and Poly(ethylene oxide) as Studied by High-Resolution Solid-State 13 C NMR Spectroscopy, European Polymer Journal, 41, 2354–2359 (2005). 38. S.W. Kuo and F.C. Chang, Studies of Miscibility Behavior and Hydrogen Bonding in Blends of Poly(vinylphenol) and Poly(vinylpyrrolidone), Macromolecules, 34, 5224–5228 (2001). 39. A. Asano, Y. Murata and T. Kurotsu, Crystallinity and Miscibility of Poly(vinyl isobutyl ether)/Poly(ε - L- Lysine) Blends by Solid State 13 C NMR Study, e-Journal of Soft Materials, 3, 1–8 (2007). 40. X. Li, W. Fu, Y. Wang, T. Chen, X. Liu, H. Lin, P. Sun, Q. Jin and D. Ding, Solid-State NMR Characterization of Unsaturated Polyester Thermoset Blends Containing PEO–PPO–PEO Block Copolymers, Polymer, 49, 2886–2897 (2008). 41. D.M. Schachter, J. Xiong and G.C. Tirol, Solid Dtate NMR Perspective of Drug–Polymer Solid Solutions: A Model System Based on Poly(ethylene oxide), International Journal of Pharmaceutics, 281, 89–101 (2004). 42. J-Z. Chen, S.V. Ranade and X-Q. Xie, NMR Characterization of Paclitaxel/poly(styrene-isobutylene-styrene) Formulations, International Journal of Pharmaceutics, 305, 129–144 (2005). 43. I. Ando, T. Kameda and N. Asakawa, Polypeptides, in Solid State NMR of Polymers, I. Ando, T. Asakura (Eds), Elsevier, Amsterdan, 819–851, 1998. 44. S. Maeda, K-K. Kunimoto, C. Sasaki, A. Kuwae and K. Hanai, Characterization of Microbial Poly (ε-L-lysine) by FT-IR, Raman and Solid State 13 C NMR Spectroscopies, J. Mol. Struc., 655, 149–155 (2003). 45. P. Hazendonk, C. de Denus, A. Iuga, P. Cahoon, B. Nilsson and D. Iuga, A Morphological Study of Poly[bis(trifluoroethoxy)phosphazene] using Solid-State NMR: Introducing Domain Selective 1 H and 19 F Decoupled 13 C MAS NMR, J. Inorg. Organomet. Polym. Mater., 16, 4 (2006). 46. P. Wormald, B. Ameduri, R.K. Harris and P. Hazendonk, Fluorine-19 Solid State NMR Study of Vinylidenefluoride Polymers using Selective Relaxation Filters, Solid-state Nucl. Mag., 30, 114 (2006). 47. P. Hazendonk, T. Montina, A. Iuga and P. Wormald, The DIVAM filter: A NMR Signal Selection Method that is Tunable to Various Structural Domains Within Semicrystalline Materials, Mater. Res. Soc., 984, (2007). 48. A.S. Borisov, P. Hazendonk and P.G. Hayes, A Morphological Study of Poly[bis(trifluoroethoxy)phosphazene] using High Resolution Solid-State 1 H, 19 F, 31 P and 13 C NMR Spectroscopy, J. Inorg. Organomet. Polym. Mat., 18, 163–174 (2008). 49. M. Trejo-Dur´an, A. Mart´ınez-Richa, R. Vera-Graziano, G. Mendoza-D´ıaz and V. Casta˜no-Meneses, Silicon-29 and Carbon-13 Nuclear Magnetic Resonance Identification of Intermediates Developed during the Formation of a Hybrid Based on TEOS and DCN, J. Appl. Polym. Sci., 99(2), 520–531 (2006).

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50. L. Dewimille, B. Bresson and L. Bokobza, Synthesis, Structure and Morphology of Poly(dimethylsiloxane) Networks Filled with in situ Generated Silica Particles, Polymer, 46, 4135–4143 (2005). 51. J. Huang, C. He, X. Liu, J. Xu, C.S.S. Tay and S.Y. Chow, Organic–Inorganic Nanocomposites from Cubic Silsesquioxane Epoxides: Direct Characterization of Interphase, and Thermomechanical Properties, Polymer, 46, 7018–7027 (2005). 52. P. Hajji, L. David, J. Gerard, J. Pascault and G. Vigier, Synthesis, Structure, and Morphology of Polymer–Silica Hybrid Nanocomposites Based on Hydroxyethyl Methacrylate, Journal of Polymer Science: Part B: Polymer Physics, 37, 3172–3187 (1999). 53. T.C. Farrar, E.D. Becker, Pulse and Fourier Transform NMR, Academic Press, New York, 1971. 54. W. Chen, H. Feng, D. He and C. Ye, High Resolution Solid-State NMR and DSC Study of Poly(ethylene glycol)silicate Hybrid Materials via Sol-Gel Process, Journal of Applied Polymer Science, 67, 139–147 (1998). 55. P. Adriaensens, R. Dams, L. Litsen D. Vanderzande and J. Gelan, Study of the Nanomorphology of OC1 C10 PVP/precursor-PPV Blends by Solid State NMR Relaxometry, Polymer, 45, 4499–4505 (2004). 56. P. Adriaensens, L. Storme, R. Carleer, J. Gelan and F.E. DuPrez, Comparative Morphological Study of Poly(dioxolane)/poly(methyl methacrylate) Segmented Networks and Blends by 13 C solid-state NMR and Thermal Analysis, Macromolecules, 35, 3965–3970 (2002). 57. S. Joseph, F. Laupretre, C. Negrell and S. Thomas, Compatibilising Action of Random and Triblock Copolymers of Poly(styrene-butadiene) in Polystyrene/polybutadiene Blends: A Study by Electron Microscopy, Solid-State NMR Spectroscopy and Mechanical Measurements, Polymer, 46, 9385–9395 (2005). 58. C. Hucher, R.-P. Eustache, F. Beaume and P. Tekely, Motional Heterogeneity in Poly(ether-block-amine) Copolymers as Revealed by Solid-State NMR, Macromolecules, 38, 9200–9209 (2005). 59. W. Lequieu, P. Van De Velde, F.E. Du Prez, P. Adriaensens, L. Storme and J. Gelan, Solid-State NMR Study of Segmented Polymer Networks: Fine-tuning of Phase Morphology via their Molecular Design, Polymer, 45, 7943–7951 (2004). 60. S. Bourbigot, X. Flambard, B. Revel, Characterisation of Poly(p-phenylenebenzobisoxazole) Fibres by Solid State NMR, European Polymer Journal, 38, 1645–1651 (2002). 61. F. Samyn, S. Bourbigot, C. Jama, S. Bellayer, S. Nazare, R. Hull, A. Castrovinci, A. Fina, G. Camino, Crossed Characterization of Polymer-Layered Silicate (PLS) Nanocomposite Morphology: TEM, X-ray Diffraction, Rheology and Solid-State Nuclear Magnetic Resonance Measurements, European Polymer Journal, 44, 1642–1653 (2008). 62. A. Domjan, G. Erdodi, M. Wilhelm, M. Neidhofer, K. Landfester, B. Ivan and H. W. Spiess, Structural Dtudies of Nanophase-Separated Poly(2-hydroxyethyl methacrylate)-1-polyisobutylene Amphiphilic Conetworks by Solid-State NMR and Small-Angle X-ray Scattering, Macromolecules, 36, 9107–9114 (2003). 63. A. Buda, D.E. Demco, M. Bertmer, B. Blumich, B. Reining, H. Keul and H. Hocker, Domain Sizes in Heterogeneous Polymers by Spin Diffusion using Ssingle-Quantum and Double-Quantum Dipolar Filters, Solid State Nucl. Magn. Reson., 24, 39–67 (2003). 64. E.R. de Azevedo, T.J. Bonagamba and D. Reichert, Molecular Dynamics in Solid Polymers, Progress in Nuclear Magnetic Resonance Spectroscopy, 47, 137–164 (2005).

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14 ESR Spectroscopy of Multiphase Polymer Systems Sre´cko Vali´c School of Medicine, University of Rijeka, Croatia, and Rudjer Boˇskovi´c Institute, Zagreb, Croatia

Mladen Andreis Rudjer Boˇskovi´c Institute, Zagreb, Croatia

Damir Klepac School of Medicine, University of Rijeka, Croatia

14.1 Introduction Electron spin resonance (ESR) spectroscopy gives the possibility to detect unpaired electrons [1–3]. This spectroscopic method is also known as electron paramagnetic resonance (EPR) spectroscopy but the former name is now more frequently used. Mostly, polymers do not give rise to an ESR spectrum since they do not possess intrinsic paramagnetism (unpaired electrons). In order to study such materials by ESR, the paramagnetic centers must be introduced into the matrix [1, 4, 5]. Nitroxyl radical has an unpaired electron that belongs to the nitroxyl group and occupies the pz orbital of the nitrogen atom. The nitroxyl group is often surrounded by four methyl groups, substituted in α-positions, which sterically protect the paramagnetic center from a possible attacking reagent and hence contribute to the stability of a free radical. The most important property of a nitroxyl radical for the study of a polymer system is that it must be compatible with the polymer matrix. Therefore, it is not enough to take care only of miscibility between the nitroxyl radical and the polymer matrix, but also of the radical size, shape, polarity and flexibility [6, 7]. A good nitroxyl radical must be able to follow the motional behavior of polymer chain segment. Two approaches to the study of polymer systems by nitroxyl radicals are commonly known: (i) radicals can be covalently attached to a polymer chains, or (ii) they can be introduced mechanically (by diffusion) into Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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Handbook of Multiphase Polymer Systems Spin probed polymer

R

N

Spin labeled polymer

O

Figure 14.1

Schematic presentation of a spin probe and spin label techniques.

a polymer matrix. The former is named spin labeling and the latter is the spin probe method (Figure 14.1). The advantage of spin labeling is the precisely known position of the nitroxyl radical in a polymer chain [7]. However, this method demands additional synthetic efforts and is rather limited since it cannot be applied to polymer systems without an adequate functional group. The spin probe method is easier and gives the possibility to introduce nitroxyl radicals into polymer systems simply by dissolving or swelling the polymer in a probe solution. After the removal of solvent, probes stay incorporated in a polymer matrix. Recent investigation indicates that the temperature of the probe incorporation process must be strictly controlled, particularly for inhomogeneous systems [5, 8]. In addition, when the spin probe method is used, it is necessary to take care of a possible attachment of the spin probe to the polymer chain by hydrogen bond. In such a case, the effect of the hydrogen bond may influence T 5mT value (defined below) [9]. In order to study the mechanism and kinetics of radical reactions in polymer systems by ESR, a spin trap technique is usually applied. The technique is based on the reaction of extremely short half-life reacting radicals with spin trap molecules resulting in a stable free radical. One of the most successful spin trapping agents is phenyl-tert-butyl-nitrone (PBN). A typical reaction of PBN with a short half-life reacting radical (R·) is shown in Figure 14.2. A big advantage of ESR spectroscopy when compared to other techniques is that very low concentrations of nitroxyl radicals (∼ 0.1 wt % or lower) embedded in a polymer matrix result in a good signal-to-noise ratio. Consequently, the introduction of nitroxyl radicals into the polymer matrix induces only a slight perturbation of the matrix with no significant influence on its intrinsic properties. The nitroxyl radicals used in the papers cited in this chapter are presented in Table 14.1.

CH

O

R O

N C(CH3)3 + R • +

C N• C(CH3)3 H

Figure 14.2

Schematic presentation of a trapping reaction of PBN.

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Table 14.1 Nitroxyl radicals used in papers cited in this chapter. Structure

R

Name (abbreviated)

(a) Piperidine derivatives based on 2,2,6,6-tetramethylpiperidine-1-oxyl R

N O

I II III IV V VI

-H =O -OH -NH2 -COC15 H31 -OCOC6 H5

TEMPO TEMPONE TEMPOL TAMINE BzONO

(b) Pyrrolidine derivatives based on 2,2,5,5-tetramethylpyrrolidine-1-oxyl R

N O

VII

-CO(OCH2 CH2 )n OH

(c) Pyrroline derivatives based on 2,2,5,5-tetramethylpyrroline-1-oxyl R

N O O N CH3

O O

CH2 CH2 O

CH2 CH O n

N

VIII

O

CH2 CH2 O m

n

O

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Table 14.2 (Continued) Structure

R

Name (abbreviated)

N

O

IX (d) Oxazolidine derivatives R1

R2

O

N

X XI XII XIII XIV XV XVI XVII

O

R1: -C13 H27 R2: -C3 H6 COOH R1: -C11 H23 R2: -C5 H10 COOH R1: -C8 H17 R2: -C8 H16 COOH R1: -C2 H5 R2: -C14 H28 COOH R1: -C11 H23 R2: -C5 H10 COOCH3 R1: -C8 H17 R2: -C8 H16 COOCH3 R1: -C5 H11 R2: -C4 H9 R1 = R2: -C9 H19

5-DSA 7-DSA 10-DSA 16-DSA 7-DSE 10-DSE 5-DD 10-DND

O N

O

XVIII (e) Diaryl derivatives H3C O

XIX

N O

O

CH3

DPAN

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This chapter is focused on the results of ESR studies of multiphase polymer systems published during last decade. Previous works in this field were systematically described elsewhere [6]. It should be remarked here that a considerable number of papers represent the ESR studies of biopolymers [10, 11] and homopolymers are not included in this overview.

14.2 Theoretical Background Once the nitroxyl radical is placed in a static magnetic field, the interaction energy of the magnetic moment associated with the spin S with an external magnetic field B is described by the Hamiltonian [1, 2]: H = μB B gS + I AS

(14.1)

where μB is the Bohr magneton and I is the nuclear spin of 14 N. The above Hamiltonian consists of two terms: first, known as the Zeeman interaction, which describes the interaction of the electron spin S with the static magnetic field B and the second, the hyperfine interaction, which describes the interaction between electron and nuclear magnetic moments. Values g and A are the tensors which parameterize the Zeeman and hyperfine interaction, respectively. The Zeeman interaction splits the energy of electron spins in two levels and the hyperfine interaction additionally splits each Zeeman level in three sublevels, as shown in Figure 14.3. Finally, the ESR spectrum of a nitroxyl radical shows three absorption lines, registered in the form of the first derivative. As long as magnetic interactions are time averaged by fast molecular motions the ESR spectrum exhibits three well-resolved Lorentzian lines of equal intensities (in the limit of isotropic motions) characteristic of the fast motions, corresponding to three projections of the nitrogen spin on the direction of the external

E

B=0

B>0 ms = +1/2

mI = +1 mI = 0 mI = –1

ΔE = h ν

ms = –1/2

mI = –1 mI = 0 mI = +1

B

Figure 14.3 Energy levels of a nitroxyl radical placed in an external magnetic field; ms is the magnetic quantum number of electron spin and mI is the nuclear spin quantum number of 14 N (nuclear spin I = 1).

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I(–1)

I(+1)

ΔB b

2Azz

In c

Im

Ib

Figure 14.4 Model ESR spectra of a nitroxyl radical typical for the polymer: (a) well above Tg ; (b) at ‘rigid limit’; and (c) inhomogeneous system just above Tg . Adapted from [8]. Copyright © 2008 by John Wiley & Sons, Inc.

magnetic field [6]. However, when motions slow down, the line widths increase and line intensities become slightly different. The model ESR spectrum of a spin probe in a polymer matrix measured well above the glass transition temperature (Tg ) is shown in Figure 14.4(a). When the segmental motions are frozen (T < Tg ), ESR spectrum indicates the slow dynamics of the spin probe determined mostly by the size of free volume holes. However, the motions of side chain groups can also contribute to the slow tumbling of probe molecules. A typical rigid limit spectrum (defined below), composed of three overlapping broad lines, is presented in Figure 14.4(b). In a particular case of an inhomogeneous system, the ESR spectrum observed in the temperature region just above Tg shows a bimodal character [6, 8], which can be seen in Figure 14.4(c). Such a bimodal spectrum is composed of two components: the broad component attributed to the slow motion and the narrow one that corresponds to the fast motion of a spin probe. The simplest analysis of a bimodal spectrum can be made by measuring the intensities of broad (Ib ) and narrow (In ) line. Then the calculated intensity ratio In /Ib can be taken as a simple measure of the probe dynamics. Higher value of this ratio indicates higher amount of the fast component present in a dynamically inhomogeneous system. When the contribution of the narrow component is relatively slow, i.e. the free volume fraction in the matrix is small, it is convenient and probable more accurate to use intensity ratios Ib and Im (central line intensity) in determining the free-volume fraction of the composite spectra [12]. The analysis of such composite spectra is often used in the characterization of different heterogeneous systems, particularly elastomers, polymer blends, copolymers, block copolymers, interpenetrating polymer networks and polymer-filler systems [6, 13–19]. Further analysis of ESR spectra makes it possible to deduce useful data about motional behavior of nitroxyl radicals, which is closely related to the dynamics of polymer matrix. Taking into consideration an isotropic Brownian motion, relatively precise values of rotational correlation times (τ R ) in the range of fast motions

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(10−11 s < τ R < 10−9 s) can be estimated using the Kivelson theory [20]: −5



τ R = 0.65 × 10 B

I (0) I (−1)

1/2



I (0) + I (+1)



1/2 −2

(14.2)

where I(–1), I(0) and I(+1) are intensities (peak-to-peak amplitudes) of low, central and high field line, respectively, and B is the line width of the central line, also shown in Figure 14.4(a). For slow spin probe motion (τ R > 10−9 s), τ R values can be calculated by: τ R = a(1 − s)b

(14.3)

where a and b are constants depending on the diffusion model, and s is the ratio of outer extrema separation (2Azz  ) of the given spectrum (Figure 14.4(b)) divided by its maximum value (2Azz ), measured for the completely immobilized nitroxyl radical (‘rigid limit’ spectrum) [21], observed at T = 77 K). On the base of calculated rotational correlation times (τ R ), it is possible to calculate the activation energy (Ea ) for relaxation processes in different temperature regions, using Arrhenius equation: τ R = τ0 e−Ea /RT

(14.4)

The preexponential factor (τ 0 ) denotes the hypothetical rotational correlation time for a free rotation of nitroxyl radical at infinite temperature, R is universal gas constant and T is thermodynamic temperature. The value 2Azz  , when measured as a function of temperature, shows a typical sigmoidal dependence, as shown in Figure 14.5. From the data presented in Figure 14.5, it is possible to determine an additional ESR parameter, known as T 5mT . This parameter corresponds to the temperature at which the outer extrema

Figure 14.5 Temperature dependence of outer extrema separation (2Azz  ) of ESR spectra for NR containing 10 phr of nanosilica and 4 phr of coupling agent bis(γ -triethoxysilylpropyl)tetrasulfide. Ta and Td are the temperatures of appearance of the narrow component and disappearance of the broad component, respectively. Adapted from [22]. Copyright 2008, with permission from Elsevier.

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separation reaches the value of 5 mT. Natural rubber (NR) is an inhomogeneous system and therefore the broad component is partially present in the ESR spectra at T > Tg . Since the decrease of 2Azz  value with the temperature increase is due to the change in motional probe behavior that reflects the dynamics of surrounding polymer chain segments, T 5mT value can be correlated with the glass transition temperature of the investigated polymer system. This correlation, based on the free volume model, is expressed by the equation [23]: T5mT − Tg = 52 {2.9 f [ln(1/ f )] − 1}

(14.5)

where f is the ratio between the activation volume of a spin probe and a polymer segment undergoing relaxation. The value of T 5mT , when compared to Tg measured by differential scanning calorimetry (DSC), is shifted toward higher temperatures. This can be explained by the large difference in the effective frequency of ESR spectroscopy, which ranges from 107 Hz to 108 Hz at T 5mT , and that of calorimetric methods [3], by which Tg values are measured at ∼1 Hz. Hence, T 5mT values of crosslinked and uncrosslinked NR samples differ for about 5 degrees, even if DSC measurements give the same T g values (–64◦ C) for both samples [22]. This demonstrates the sensitivity of the spin probe on local segmental motions. On the assumption that the nitroxyl radicals are distributed homogeneously in a dynamically inhomogeneous polymer matrix, it is possible to estimate the exact amount of the mobile and immobile matrix fraction at temperatures in the vicinity of T 5mT and above T 5mT . In order to obtain such information, it is necessary to make a decomposition of bimodal ESR spectra. This implies the data retrieval of broad and narrow components contained in bimodal ESR spectrum. In fact, spectral decomposition comprises the detection of such broad and narrow components whose superposition (a simple addition of their corresponding intensities) will result in an experimentally observed spectral line shape. Programs for calculating ESR spectra, developed by Schneider and Freed [24], Budil et al. [25] and by Stoll and Schweiger [26, 27], are available and they are often used satisfactorily taking into consideration an isotropic Brownian motion [28]. The simulation of spectra using the program developed by Freed includes the following steps: (a) simulation of the rigid limit powder spectrum of the nitroxyl radical for obtaining the principal values of the g and A tensors; (b) choice of the diffusion tilt angles (the Euler angles that specify the rotation taking the magnetic axes into those of the rotational diffusion tensor); (c) choice of the motional model; and (d) simulation using the rotational frequency as a fitting parameter (Figure 14.6). One of the motional models frequently applied for the spin label technique is microscopic order with macroscopic disorder (MOMD). This model is based on the assumption that the polymer sample consists of locally ordered domains that are randomly oriented on macroscopic scale, as presented schematically in Figure 14.7. In fact, nitroxyl undergoes microscopic molecular ordering defined by local director. In many polymer systems local directors are rigidly fixed and randomly oriented with respect to the magnetic field B [4]. However, in some inhomogeneous systems the programs have failed to reproduce experimentally observed spectra [6, 30]. Particularly, the main problem in simulating ESR spectra occurred when reproducing the spectra of slow motion. Consequently, a semiempirical method has been developed and applied for the simulation of bimodal spectra of NR [31]. This method is based on the calculation of fast motional components and their combination with a set of experimentally measured slow motional components of the like systems. One of the most useful ways in which magnetic resonance may be applied to anisotropic media is determination of the molecular order parameters: S=

 1 3 cos2 β − 1 2

(14.6)

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335

340

559

345

B / mT

Figure 14.6 ESR spectrum of nitroxyl radical III used as a spin label, bonded on the surface of SiO2 particles and incorporated in styrene-butadiene rubber, recorded at 363 K. Simulations were made assuming Brownian rotational diffusion model. Dashed lines: computed; continuous lines: experimental. Reprinted from [63]. Copyright (2004) with permission from American Chemical Society.

where β is the angle between the preferred spatial direction in the medium, usually referred to the local director and the symmetry axis of the molecule [32]. The physical background of ESR spectroscopy of nitroxyl radicals given above is focused on the continuous wave (CW) ESR technique, based on the application of microwave irradiation field of constant frequency ν and sweeping the external magnetic field B until the resonance condition is fulfilled. Polymers are usually characterized by slow molecular motions and, therefore, no complete time averaging of magnetic interactions occurs. Consequently, the conventional CW ESR technique is not sensitive enough to study such systems at lower temperatures (T < Tg ).

Figure 14.7 Two dimensional schematic presentation of local microdomains ordering and their random orientations with respect to magnetic field B. Short arrows denotes local directors. Reprinted from [39]. Copyright (2004) with permission from American Chemical Society.

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Figure 14.8 Schematic presentation of CW and pulse ESR techniques. Reprinted from [90]. Copyright (2006) with permission from Springer-Verlag.

Pulsed ESR spectroscopy based on electron spin echo (ESE) is highly sensitive and gives a possibility to measure the changes in molecular motion in the low temperature range (under Tg ) where the relaxation times are about 10−6 s. Therefore, ESE is well suited for the study of changes in nanostructure of matter (range ∼3 ± 0.1 nm) which cannot be investigated by other methods. Typical microwave pulse durations in ESE spectrometers are 10–100 ns. Pulsed ESR spectra are recorded by exciting a large frequency range simultaneously with a single high-power microwave pulse of given frequency ν at a constant magnetic field B (Figure 14.8). The electron spin echo envelope modulation (ESEEM) through electron - nuclear interactions represents a convenient tool of structural studies in the nearest environment of spin labels; it can be used for identification of the types of the nuclei in the nearest environment and determination of their spatial arrangement relative to the unpaired electron. Recently, this method has found increasing application in studies of the local environment of spin labels. The physical background of ESEEM has been systematically described in reference [33]. The dipole–dipole interaction between radicals is now widely used in studies of the nanostructure of spin labeled macromolecular systems. The pulsed double electron–electron resonance (PELDOR), also known as DEER, allows one to determine the distances between radicals on the nanometer scale, get information on the conformations of doubly spin labeled macromolecules and estimate the size of the aggregate of spin labeled molecules. In the PELDOR experiments the echo signal appears upon application of a conventional two-pulse sequence and an additional 180◦ C pump pulse at the other resonance frequency. There are also other sequences such as 3-pulse or 4-pulse PELDOR (Figure 14.9).

14.3 Copolymers Copolymers are polymer chains constituted of two or more different repeating units. Desirable chemical and physical properties of such materials can be achieved by changing the relative molar ratio of various

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Figure 14.9 Pulse sequences used in the PELDOR method. (a) Three-pulse PELDOR technique. The spin echo signal V(τ ) appears after application of the π /2 and π pulses at the frequency ν A . The pumping π -pulse applied at instant T at the frequency ν B changes the echo signal amplitude. These changes are detected as the PELDOR signal V(T). (b) Four-pulse PELDOR technique. The echo signal is induced by applying three pulses (π /2, π , π ) at the frequency ν A . Changes in the echo signal (PELDOR effect) appear after application of the pumping π pulse at the frequency ν B during the time interval from τ to τ + τ 1 . Reprinted with permission from [34]. Copyright 2008 Turpion-Moscow Ltd.

repeating units. Especially interesting are block copolymers which are able to form domains of different structures under the strong segregation limit (SSL) condition [35]. For these reasons, copolymers attract a lot of scientific attention and are the subject of numerous studies. The self-assembly of amphiphilic triblock copolymers (poly(ethylene oxide)-poly(propylene oxide)-poly(ethylene oxide)) (PEO-PPO-PEO) was investigated in dispersions of single-walled and multi-walled carbon nanotubes (SWNT and MWNT, respectively) as a function of temperature. Three labeled Pluronics (nitroxyl radical VIII with different m and n), where a spin label was covalently attached at the end of each PEO chain, were employed and therefore the motion of the PEO tails was monitored. The two spin labeled Pluronics differing in the lengths of the PEO and PPO blocks (L62-NO with m = 30 and n = 5 and P123-NO with m = 70 and n = 20) are located at different locations as presented schematically in Figure 14.10. It was found that SWNT and MWNT modify the temperature, enthalpy, and dynamic behavior of polymer self-assembly. In particular, SWNT were found to increase the cooperativity of aggregating chains and

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corona (PEO) core (PPO)

L62-NO P123-NO

Figure 14.10 Schematic presentation of locations of spin labeled Pluronics L62-NO and P123-NO in dispersions of carbon nanotubes. Reprinted with permission from [36]. Copyright 2008 American Chemical Society.

dominate aggregate dynamics. MWNT reduced the cooperativity, while colloidal carbon black additives did not show similar effects [36]. Aqueous solutions of the triblock PEO13 -PPO30 -PEO13 (Pluronic L64) were investigated over a wide concentration range (20–100%, (w/w) polymer) in the micellar, liquid crystalline and reverse micellar phases, using the spin probe method. Three amphiphilic nitroxyl spin probes based on N-doxyl-stearic acid were used to measure the local polarity, dynamics and degree of orientation in the self-assembled system. The spectra reflect hydrophobic sites where the hydration is very low (nitroxyl radical X) and where the hydration is essentially zero (nitroxyl radicals XII and XV, respectively). Simulations of the ESR spectra in the lamellar phase reflect an order parameter S, which decreases gradually from the EO/PO interface into the PO domains. The results of the present study provide evidence for the existence of PO domains with zero effective local hydration. Different values of order parameter, observed for different spin probes suggest an existence of a gradient of order in the aggregates [37]. In fact, the order parameter decreases from corona to the hydrophobic interior. The hexagonal, silica-based, mesoporous material SBA-15 prepared using PEO20 -PPO70 -PEO20 (Pluronic P123) as a template and tetramethoxy orthosilane (TMOS) as a silica source was investigated by ESR in combination with electron spin-echo envelope modulation (ESEEM); the reaction was conducted in D2 O [38, 39]. Pluronic spin probes with different PEO and PPO chain lengths were used (L62-NO with m = 30 and n = 5, P123-NO with m = 70 and n = 20 and F127-NO with m = 70 and n = 106). The spin label was located at the end of the PEO chain, which placed the probes at different regions of the micelles. According to Figure 14.11, the spectrum of L62-NO is a superposition of a mobile species, attributed to the nitroxyl radicals in the mesopore region, and an immobilized species, assigned to spin labels trapped in the micropores. The spectra of the other two spin probes show that the longer the PEO chain, the higher the relative amount of the immobilized species. By comparing the results obtained with pulse ESEEM sequence with the in situ ESR measurements of different spin probes, four time-dependent stages in the formation of SBA-15 indicating that the polymerization propagates outward from the core/corona interface were detected (Figure 14.12). The initial steps of the formation mechanism of SBA-15, focusing on the process that occurs within the individual micelles, not considering changes in shape/length are presented in Figure 14.12. Initially, the polymer chains are packed together and form micelles. After adding acid and TMOS to the micellar solution, TMOS and partially hydrolyzed TMOS, which are hydrophobic molecules, penetrate into the core

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Figure 14.11 ESR spectra of L62-NO, P123-NO and F127-NO in SBA-15, recorded at room temperature. Adapted with permission from [39]. Copyright 2004 American Chemical Society.

(stage (a)). Hydrolysis occurs very fast and hydrolyzed monomers diffuse into the corona region, where the polymerization process of the silica begins, causing a depletion of D2 O/OD from the corona region and/or migration of forming silica network outward. During the first hour (stage (b)) of the reaction, this is felt at the core/corona interface and at the second hour (stage (c)) of the reaction in the corona region. It is possible that stages (b) and/or (c) are associated with the formation of rod-like micelles, micelle aggregation, and

Figure 14.12 Proposed model for the initial stages in the formation of SBA-15 (Si/P123 = 59). Gray scale in the corona region represents the amount of D2 O/OD; dark gray denotes a large amount of water in the corona region, whereas light gray denotes a region with less water. Reprinted with permission from [39]. Copyright 2004 American Chemical Society.

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the formation of hexagonal order. After the hexagonal order is formed, there is a continuous increase of the hydrophobic core and a decrease in the silica wall thickness, as described in the model at stages (b)–(d). [39] Another group of authors used nitroxyl radicals X and XIII to determine the micropolarity and microviscosity of self-assembling systems of mmePEG-P(CL-TMC) based on monomethoxylated polyethylene glycol (mmePEG), caprolactone (CL) and trimethylene carbonate (TMC) having different PEG chain lengths and different CL/TMC ratios. It was found that the micelles with mmePEG-P(CL-TMC) polymers are biphasic in nature. The hyperfine splitting constant and the microviscosity determined from the spectra of nitroxyl radicals X and XIII were very different. The micelles are made of a more hydrophilic and less viscous corona and a more hydrophobic and more viscous core. It seems that the outer shell could be made of PEG chains, the hydrophobic part of the chains making the core. The polarity in the vicinity of nitroxyl radical X was rather stable whereas the polarity in the vicinity of nitroxyl radical XIII significantly increased with temperature [40]. Other systems made of PEG, monooleyl glycerol (MOG) and succinic anhydride (SA), with different PEG chain lengths and different components ratio were also studied. Due to the spatial configuration of random polymeric systems no clear domain separation was found. Polystyrene-poly(methyl acrylate) (PS-PMA) diblock labeled with nitroxyl radical IV at the junction points was investigated. An existence of the heterogeneous mixture of each segment in the interfacial region at a certain length scale was found. It was concluded that the dynamic heterogeneities in the interfacial region were due to the gradient of the segmental concentration in the interfacial region, and the dynamic environment in the interfacial region was strongly influenced by the mobility of the block chains within microdomains [41]. Molecular mobility in the interfacial region of a microphase-separated structure was studied in binary mixtures of AB-type diblock copolymers and homopolymers (miscible with only A). Two types of binary blends containing PS-PMA were studied: (a) with short homopolymer chains of poly(cyclohexyl acrylate) (PCHA-S) and (b) with long homopolymer chains of poly(cyclohexyl acrylate) (PCHA-L). It was shown that the mobility of the PCHA homopolymers is obviously depressed by blending with the PS-PMA since the PCHA homopolymers were localized in rigid regions (PS phase). It was also considered that almost no PCHA-L penetrated into the interfacial region. Composition in the interfacial region changed gradually from the PS phase to the PMA phase. Additionally, the study of binary mixtures of the PS-PMA and short homopolymer PS chains (PS-1) indicated that the mobility in the interfacial region was somewhat enhanced by the addition of PS-1 due to their penetration into the interfacial region. The results suggested that the spatial distribution of PS-1 in the PS phase was not completely uniform even if the molecular weight of PS-1 was small [42]. The dynamic heterogeneity of the nitroxyl radical XVIII dissolved in nematic azobenzene polymethacrylates (PMA4) homopolymer and copolymer of PMA4 and poly(methyl methacrylate) (PMMA) was studied by Andreozzi et al. Spectra were measured in the temperature range from well above nematic isotropic transition (T NI ) down to below Tg . The distribution functions of molecular sites were determined, and their stability was monitored as a function of thermal history. The different molecular site distributions and cooperativity degrees were compared and discussed [43]. An extended study of same samples has shown that the copolymer 40/60 (containing 40% of MMA) easily exhibits stable substrates without thermal treatment, necessary in the 30/70 copolymer and the corresponding PMA4 homopolymer [44]. Additionally, the time evolution of line shape in the spectra was monitored for the same samples annealed at different temperatures (T = 358 K or 383 K). It was shown that the annealing of the copolymer at temperatures in the isotropic phase favors formation of homogeneous nematic polymeric substrates. Such behavior was found to be in marked contrast with that of the homopolymer, for which a homogeneous substrate was observed in a limited temperature region after annealing at 358 K. This effect was attributed to the greater mobility of the side groups due to the presence of the methyl methacrylate units, which weaken the nematic potential [45]. The spin probe method was further used to investigate the molecular motions on nanometer scale in an oil-extended symmetric triblock of polystyrene-polybutadiene-polystyrene (PS-PB-PS) with molecular mass

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of polybutadiene block about 80000, containing 34.5% of styrene. Spectra of SBS containing various amounts of oil (1.45–28.60 wt%) doped with nitroxyl radical I as a spin probe were measured in a temperature range from 183 K to 348 K. It was concluded that at oil concentration of about 17%, a transition from bicontinuous morphology to a segregated morphology can be noted. For higher oil concentrations, a decrease of Tg of PS blocks was observed. This was explained by the diffusion of a low amount of oil molecules or by the penetration of some lubricated PB chains within the PS blocks [46]. Spectra of methanolic solutions of the copolymer of 2-hydroxyethyl methacrylate (HEMA) with spin labeled methacrylic acid copolymer at concentrations between 2 and 50 wt% were measured over a broad temperature range. Temperature dependence of the RS parameter characterizing the local segmental dynamics in PHEMA and local dynamics and conformation of the copolymer were determined by fitting theoretical to the experimental spectra calculated using the MOMD model. A conformational transition from a less compact to more compact conformation of the polymer chain of PHEMA was observed. Intramolecular interactions in a dilute solution in a certain temperature range, where PHEMA exists in a compact conformation, hinder local polymer dynamics [47]. Various potentially amphiphilic gel-type resins were prepared by copolymerization with γ -irradiation of hydroxyethyl methacrylate (HEMA), hydroxypropyl methacrylate (HPMA) and trimethylolpropane trimethacrylate (TMPTM); the nitroxyl radical II was used as a spin probe. The materials produced the expected swelling behavior when a nonpolar liquid (THF) was chosen as swelling medium. Although HEMA is a hydrophilic monomer, polymer sequences rich in HEMA may mutually interact very strongly via hydrogen bonding to make their solvation by water scarcely beneficial in terms of the overall swelling of the polymer framework. Accordingly, the swelling of poly-HEMA itself in water was seen to be definitely modest [48]. Microphase separation in poly(acrylonitrile-butadiene-styrene) (ABS) was studied as a function of the butadiene content and method of preparation using nitroxyl radicals XVI and XVII as spin probes. Results for the ABS copolymers were evaluated by comparison with similar studies of the homopolymers polybutadiene (PB), polystyrene (PS), and polyacrylonitrile (PAN) and the copolymers poly(styrene-co-acrylonitrile) (SAN) and poly(styrene-co-butadiene) (SB). The spectra were measured in the temperature range 120 K–420 K and were analyzed in terms of line shapes, line widths, and hyperfine splitting; the appearance of more than one spectral component was taken as an indication of microphase separation. In the spectra of nitroxyl radical XVII in all homopolymers, as well as in the copolymers SAN and SB, only one motional component was detected. In contrast, two spectral components differing in their dynamic properties were detected for both probes in the three types of ABS samples studied and were assigned to spin probes located in butadiene-rich domains (the fast component) and SAN-rich domains (the slow component) (Figure 14.13) [49, 50]. The simulations of two-component ESR spectra showed that the jump diffusion model (the spin probe has a fixed orientation for a given time and then jumps instantaneously to a new orientation) leads to good agreement with experimental spectra (Figure 14.14). The parameters used to calculate the spectra are given in reference [50]. Nitroxyl radicals formed in thermally-treated heterophasic propylene-ethylene copolymers (HPEC) containing bis(2,2,6,6-tetramethyl-4-piperidinyl) sebacate (Tinuvin 770) as a hindered amine stabilizer (HAS) were studied in the temperature range from 100 K to 433 K. The spin probes composed of a mixture of monoand bi-radicals are derived from the HAS and are termed HAS-NO. This study has demonstrated the exceptional sensitivity of ESR spectra from nitroxyl radicals to polymer morphology and degree of crystallinity. The results of thermal ageing studies suggested a faster degradation rate in HPEC containing a higher amount of ethylene component [51, 52]. A similar study, including also the ESR imaging techniques, was focused on thermal and UV oxidative degradation of ABS copolymer containing HAS. Upon UV irradiation provided by fluorescent lamp (290 nm – 320 nm), the nitroxyl signal was initially strong only on the irradiated side and increased with time through the sample. These results were taken as evidence for extensive damage

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Figure 14.13 ESR spectra of (A) nitroxyl radical XVII (denoted as 10DND) in ABS1, ABS2, and ABS3 at 330 K and (B) nitroxyl radical XVI (denoted as 5DD) in ABS1, ABS2, and ABS3 at 320 K (ABS1, ≈10 % B), (ABS2, ≈25 % B), (ABS3, ≈45 % B). The 2Azz  of the slow component is given on the left. Adapted from [49]. Copyright © 2002 by John Wiley & Sons, Inc.

Figure 14.14 Experimental (solid lines) and calculated (dashed lines) ESR spectra of ABS2 copolymer doped with nitroxyl radical XVII for the indicated temperatures. τ R values are given on the left and relative intensity of the narrow/fast component (%F) is deduced from the simulations. Adapted from [50]. Copyright © 2002 by John Wiley & Sons, Inc.

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on the irradiated side and slower thermal degradation in the entire sample. To the contrary, after irradiation with the Xe source (closely resembles the solar spectrum), the nitroxyl radicals were detected at both the irradiated side and the opposite side, and their intensity was negligible in the sample interior. With both UV sources a spatial variation of the ESR line shapes was observed. There was no spatial variation of the ESR line shapes during thermal degradation at 318 K and 333 K and the spatial distribution of the nitroxyl radicals was homogeneous. This study suggested a hierarchical variation of the HAS-derived radical concentration within morphological domains in ABS on the scale of a few micrometers and within the sample depth on the scale of millimeters [53]. Poly(styrene-methacrylic acid) random copolymer, containing spin labeled methacrylate units based on selectively deuterated nitroxyl radical IV distributed statistically along the polymer main chain was studied in order to investigate the solvent dependence of polystyrene local segmental dynamics in dilute solutions. The data show that the copolymer in various solvents (good, Θ and marginal solvent) exhibit non-Kramers behavior characterized by the potential barrier for local conformational transitions (Ea ) about 10.5 kJ/mol [54]. ˇ Culin et al. studied the inhomogeneity of segmental dynamics in polymer mixtures of segmented polyether polyurethane (PU) with carboxylic groups and methacrylic copolymer (PM) with ternary amine groups by spin probe and spin label methods using nitroxyl radical IV. Labeled samples were prepared by amideester interchange on PM component. Spin probed PU/PM mixtures indicated that the probe motion and the phase separation deduced from the temperature-dependent ESR spectra are sensitive to a free volume determined by the polymer–polymer interactions. It was shown that the generation of free volume is slower in the mixtures with higher functional groups content, i.e. with increased polymer–polymer interactions [55]. The motional heterogeneity of polystyrene-polybutadiene (PS-PB) diblock was studied by the spin probe technique using nitroxyl radicals III and VI. Two spectral components were observed for both spin probes above the glass transition temperature of polybutadiene blocks. The broad and narrow spectral components were attributed to polystyrene and polybutadiene domains, respectively. It was found that the range with two spectral components (the temperature interval between Ta and Td ) and T 5mT value depend on the block length. The ESR spectra of PS-PB copolymers casted from different solvents have shown that the motion of probe III is sensitive to the domain morphology created by the solvent [56]. The segmented PU based on isophoronediisocyanate, polycaprolactone, and 1,4-butanediol with different amounts of functional groups introduced into the hard segments via second chain extender, 2,2 -bis-(hydroxymethyl) propionic acid, were investigated by nitroxyl radical IV used as a spin label. Spectra of functionalized PU consisting of fast and slow component were observed. This suggests that PU hard segments are partitioned in two motionally different environments. It was concluded that motional heterogeneity increases with an increase of functional groups up to 0.35 mmol/g and above this concentration slow component decreases indicating higher degree of phase mixing and stronger effect of soft segments [57]. Temperature dependence ESR spectra of 1,1-diphenyl-2-picrylhydrazyl (DPPH) radical solubilized in aqueous solutions of aggregates of amphiphilic block copolymers with hydrophobic dendritic pendants was investigated. The results showed that the radicals exhibit a single-line ESR absorption, which is narrowed by the interspin interaction, and indicates the assembly formation of DPPH radicals in polymer aggregates [58]. The ion clusters α,ω-macrozwitterionic PS-PI, prepared by anionic polymerization and subsequent introduction of the ionic end groups, and μ,ω-macrozwitterionic and -monoionic PS-PI, prepared by anionic polymerization using 1-(4 -(N,N-dimethylamino)methylphenyl)-1-phenylethene (MDPE) as monomer between the blocks and subsequent quaternization of the tertiary amino group, were studied by variable temperature CW ESR and DEER in order to characterize the size of ion clusters, distances between them, and the dynamics of their environment in ionically functionalized diblock copolymers. 4-carboxy-TEMPO converted in its potassium salt (K-T) was used as a spin probe. Variable temperature CW ESR measurements reveal a bimodal distribution of the rotational correlation times of ionic spin probes attached to ion clusters in μ,ω-macrozwitterionic and monoionic block copolymers if the ion clusters are situated in the block

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Figure 14.15 Temperature dependence of T2 and T1 for the probe K-T in PS-PI copolymer. Adapted from [60]. Copyright 2002, with permission from Elsevier.

copolymer interface. The slow and fast component of spectra were attributed to spin probes on the ion cluster surface exposed to the more rigid PS and the more mobile PI microphase, respectively. In the case of the μ,ω-macrozwitterionic sample the fraction of cluster interface with PI was increased by a factor of 2 with respect to that of the monoionic species with a charge in the interface. DEER measurements showed no significant difference in the cluster size or intercluster distance for comparable block copolymers with ionic groups either at the block junction point or at the chain end. For α,ω-macrozwitterionic block copolymers, ion cluster sizes were found to be independent of the molecular mass of the ionomers between 10000 and 50000. It was also found that the intercluster distance does not depend on molecular mass [59]. Polystyrene-polyisporene diblock copolymers (PS-PI) probed with nitroxyl radicals I and K-T were studied with saturation-recovery (SR), ESE and stimulated echo (SE) techniques to measure electron spin-lattice (T 1 ) electron spin–spin relaxation (T 2 ) and T SE time, respectively (Figure 14.15). From the SR and SE experiments it was concluded that the reorientation of K-T attached to the ionic cluster is clearly uniaxial. This may be appreciated by the approximate equalities T 1x ≈ T 1y < T 1z and T SE x ≈ T SE y < T SE z . As the relaxation times T 1 and T SE are affected by the dynamics of K-T, it was concluded that the reorientation of the x and y axis of the molecular-frame are nearly equivalent and faster than the one of the z

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axis. At the same time T 2x = T 2y < T 2z , while on cooling from room temperature T 2 first increases and then decreases by lowering the temperature below about 200 K. It was noted that at low temperatures T 2 is mainly affected by the thermally activated rotation of the methyl groups in the nitroxyl moiety which modulates the electron-nuclear dipolar interaction with the methyl-group protons. For T > 200 K the effect is averaged out and T 2 is set by the reorientation process of the spin probe [60].

14.4 Grafted Polymers Thermal properties of copolymers, poly(N-isopropyl acrylamide), PNIPA, grafted with different amounts of PEO, were investigated in aqueous solutions using nitroxyl radical X as a spin probe because of its capability of dissolving in hydrophobic structures. The results revealed that a small change in the number of PEO chains, and probably also in their distribution along the PNIPA chain, affects the binding of the probe to the polymer. The number of PEO chains changes the size of the hydrophobic core; the polar head of the probe including the nitroxyl moiety resided outside the hydrophobic part of the polymer in a polar surrounding composed of PEO chains and water [61]. Some grafted polymers were studied with nitroxyl radicals X and XIII as spin probes with two ways of sample preparation that resulted in a different localization of spin probes in the polymers. Since the probe resided either in the outer PEO shell close to the aqueous phase or in a less polar phase close to the PNIPA core of the polymer, no motional heterogeneity was observed. The rate of the rotational diffusion of the probe increased discontinuously during the polymer collapse in the former case but in the latter, a typical linear activation of the thermal motion was observed. The results of fluorescence and ESR spectroscopies suggest that the distribution of the PEO grafts along the PNIPA chains depends on the conformation of the functionalized PNIPA during the grafting, and this affects the behavior of the graft copolymers [62]. Methacrylate (MA) monomers were graft copolymerized on an isotactic polypropylene (PP) film surface by the UV-irradiation method. Tethered PMMA and poly(butyl methacrylate) (PBMA) chains were spin labeled with nitroxyl radical IV to study molecular motion of the PMMA and PBMA chains in toluene by ESR. The spectra were composed of two components arising from two kinds of spin labels, A- and B-labels. They were a mobile label (A) in isolated single-chain globules and smaller pinned micelles and a rigid label (B) in larger pinned micelles and the gel phase on the PP surface. The fractional amount of the mobile labels calculated from spectral intensities of the decomposed spectra decreases with increasing the grafting amount and degree of orientation of the PP chains (Figure 14.16). It was also found that the PMMA chains were more mobile than the PBMA chains in the isolated single-chain globules on the PP surface despite the higher glass transition temperature of the PMMA in the bulk than that of PBMA [63]. Molecular motion of the grafted poly(methacrylic acid) (PMAA) onto the PTFE (PTFE-g-PMAA) was studied by labeling the PMAA carboxyl group with nitroxyl radical IV. Since the Arrhenius plot of τ R vs. 1/T exhibits three regions with two turning points at about 295 K and 385 K, which may correspond to the transition points for PTFE conformational changes at 292 K and 395 K, respectively, it was suggested that the molecular motion of the grafted PMAA chain was mainly controlled by that of the principal PTFE chain. The evaluated activation energy was lower for the grafted PMAA chain (e.g. 6.5 kJ/mol at temperature range below 295 K) than that for the PMAA homopolymer (13.2 kJ/mol) [64].

14.5 Blends Blends of poly(propylene carbonate) (PPC) with copolymer poly(styrene-co-4-vinyl phenol) (STVPh) have been studied with nitroxyl radical III as a spin probe. The results indicated that the nitroxyl radical existed in both, PPC-rich and STVPh-rich micro domains, corresponding to the fast motion and slow motion

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Figure 14.16 Dependency of fraction of mobile component of ESR spectra on grafting amount (GA). Mobile fraction of PMMA grafted on the PP (solid circle), PMMA grafted on the stretched PP (open circle) by photoactivating, and PMMA grafted on the PP by thermal activating (open square). Reprinted from [63]. Copyright 2004 American Chemical Society.

components in the spectra, respectively. In order to acquire quantitative information about the relative fraction of the fast motion component in the composite spectra, a simplified estimate of the fast motion content can be obtained as shown in Figure 14.4(c). The fast motion fraction increased with increasing the hydroxyl group content in copolymer STVPh. Moreover, the ESR parameters T 5mT and τ R , as well as activation energies (Ea ) showed similar dependence on the hydroxyl group content as the fast motion fraction. It resulted from the enhancement of the hydrogen-bonding interaction between the hydroxyl groups in STVPh and the carboxyl groups and ether oxygen in PPC [65]. Different hydroxyl content poly(styrene-co-(hexafluoro-2-hydroxylisopropyl)-α-methylstyene) [PS(OH)-X] copolymers were synthesized and blends with nitroxyl radical III end spin labeled PEO were studied. Two spectral components with different rates of motion were observed in the composite spectra of all the blends, indicating the existence of microheterogeneity at the molecular level. According to the variations of ESR spectral parameters Ta , Td , T, T 5mT and τ R , with the increasing hydroxyl content in blends, it was shown that the extent of miscibility was progressively enhanced due to the controllable hydrogen bonding interaction between the hydroxyl in PS(OH) and the ether oxygen in PEO [66]. The similar behavior of two component spectra is observed in linear polymer blends and semi-interpenetrating polymer networks based on poly(styrene-co-methacrylic acid) (STMAA) and poly(butyl methacrylate) (PBMA) using the same nitroxyl radical as a spin probe [67]. The PS(OH) copolymers were blended with poly(propylene carbonate) [PPC] and their miscibility, microheterogeneity and hydrogen bonding interactions were investigated with three spin probes. Nitroxyl radical I is not sensitive to the investigated system (spectra do not reveal information on micro phase separation or miscibility), while in the spectra of blends probed with nitroxyl radicals III and IV two spectral components were observed in different temperature ranges arising from the more mobile PPC-rich micro phase and the more rigid PS(OH)-rich micro phase. The τ R depends on the polymer matrix rigidity and the strength of the hydrogen bonding between the probe molecule and polymer matrix [68]. Local segmental dynamics of poly(cyclohexyl methacrylate) (PCHMA) and poly(cyclohexyl acrylate) (PCHA) in their miscible blend was studied by labeling the PCHMA and PCHA chains with nitroxyl radical

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IV at chain ends or inside sites, respectively. Due to the sensitivity of ESR techniques to molecular dynamics in multiphase systems a specific chain dynamics of the blend components can be detected separately even though a single glass transition was observed by calorimetric measurements [69].

14.6 Crosslinked Polymers The change of network morphology of various unfilled sulfur crosslinked NR samples exposed to thermal oxidative ageing has been studied with nitroxyl radical III as a spin probe. By combining the local network density distribution deduced from the complex ESR spectra and the equilibrium swelling density, it is shown that the major changes occur in the networks with predominantly polysulfidic bonds (conventional network), while the effect of ageing on the network with predominantly monosulfidic bonds is much less pronounced [70]. Two types of NR films crosslinked by γ -irradiation were investigated: (i) samples irradiated in the relaxed state and (ii) samples irradiated in the state of uniaxial deformation. It has been shown that the spectra of nitroxyl radical III diffusing in the NR matrix crosslinked under deformation are influenced by both, the local dynamics of the chain segments and their spacious orientation. The application of deformation during the crosslinking leads to the lower crosslink density. Such an effect can be attributed to the loss of the gel component induced mechanically [22]. Crosslinked polyurea probed with nitroxyl radical II dissolved in various swelling media qualitatively confirms the presence of nanoporosity. However, the spectra reveal that the spin probe undergoes a unique motion regime in each accessible nanopore, and there is no indication of nanodomains in which the probe rotation is either strongly slowed down or even hindered. It should be noticed that probe dynamics was regularly faster in unconfined solvents than inside swollen nanoporous domains, with an exception of water [71]. Besides chemically crosslinked polymers (covalent network), a number of papers deal with physically crosslinked polyelectrolyte networks as well as spin probed polyelectrolytes and ionic/nonionic surfactants in solution [72–79]. However, these studies are beyond the scope of this review.

14.7 Semi-Interpenetrating Networks (SIPNs) The motional transition and heterogeneity of SIPNs based on segmented polyester polyurethane (PU) with carboxylic groups and methacrylic copolymer (PM) with tertiary amine groups as a function of concentration of functional groups in both prepolymers were studied with nitroxyl radical I as a spin probe. The temperature-dependent spectra are sensitive to polymer interactions imposed by functional groups which determine the free volume distribution in the matrix and temperature at which motional transition takes place. The fraction of free volume (Ib /Im , see Figure 14.4(c)) increases with functional group concentration and reaches its maximum at 0.25 mmol g−1 while further increases in the functional group concentration reduce the free volume. These results are discussed in terms of the interference of plasticizing effect of the PU component and the formation of possible cluster crosslinks, which restricts segmental motions [80]. The same SIPNs with different functional group concentrations (SIPN-0, SIPN-25, SIPN-35 and SIPN-45; numbers indicating the concentration of functional groups in mmol g−1 ) as well as corresponding polymer mixtures (1:1 mass ratio) were studied by the samples containing PM component spin labeled with nitroxyl radical IV. They were used to characterize the heterogeneity of segmental motion and transitions due to the additional polymer interactions imposed by complementary functional groups, and the results were deduced from the temperature-dependent spectra. Two component spectra reflect the effect of PU chains on segmental motion

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Figure 14.17 ESR spectra of spin labeled PM in SIPN-0 measured at different temperatures. Two spectral components are marked with b (broad) and n (narrow). Adapted from [81]. Copyright 2005, with permission from Elsevier.

of the PM component below the macroscopic Tg and the ratio In /Ib is related to the complex polymer-polymer interaction or extent of miscibility (Figure 14.17). The dependence of molecular dynamics on the functional group concentration using spin label is similar to that observed with spin probe. PU/PM mixtures reveal similar motional behavior as SIPNs of the same composition. However, the differences in the fractional amount of fast and slow motions confirm better interpenetration and interaction of the two polymers in the SIPNs. Moreover, the largest fraction of fast component was found in the samples with no functional groups [81]. If the same SIPNs and mixtures are spin labeled with nitroxyl radical IV at PU component, the critical concentration of functional groups of 0.35 mmol g−1 (Figure 14.18), above which the motional restriction decreases, is observed. This is discussed in terms of the change of local packing density. Note that the motional behavior of SIPN-0 is different in comparison with three other samples containing functional groups, whose spectra above T 5mT consist of two components [17].

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Figure 14.18 Temperature dependence of 2Azz  for SIPNs of spin labeled PU with various functional group concentrations: SIPN-0 ( d), SIPN-25 (), SIPN-35 () and SIPN-45 (♦). The full points indicate the slow components coexisting with the fast ones. Adapted from [17]. Copyright 2005, with permission from Elsevier.

The morphology and miscibility SIPNs based on poly(styrene-co-methacrylic acid) [P(S-co-MAA)] of different carboxylic acid contents and poly(ε-caprolactone) (PCL) have been studied using nitroxyl radical III as a spin label. While the ESR spectra of spin labeled PCL exhibits one motional component at any specific temperature, two spectral components are observed for all SIPNs over a temperature range T (T = Td − Ta ) arising from two microphases: a PCL-rich microdomain and a P(S-co-MAA)-rich microdomain. The miscibility can be improved by increasing the carboxylic acid content, which could enhance the hydrogenbonding interactions between the ester groups of PCL and carboxylic acid groups in P(S-co-MAA). It was also found that the microphase separation became more remarkable with increasing intracomponent crosslinking of the SIPNs [82]. A series of SIPNs from castor-oil-based polyurethane (PU) and benzyl starch (BS) was investigated with nitroxyl radical III as a spin probe and spin label (dispersed in SIPN matrices and covalently bonded to PU chains). The spectra of both spin probed and spin labeled SIPNs, consisting of a single motional component, confirmed that the segment volume of PU in the SIPNs increased with the addition of BS, i.e. the chain stiffness increased as a result of strong interactions between PU and BS macromolecules [83].

14.8 Composites The diffusion of light hydrocarbons in composite polymers based on high impact polystyrene (HIPS) and polybutadiene (PB) was studied with nitroxyl radical I as a spin probe. Spin probes were dissolved in cyclopentane and the samples were dipped into the solution. Two component spectra were deconvoluted and the molar percentage of spin probe inside the PB phase (narrow component) was calculated. The finding that nearly 20% of probe molecules penetrated into a composite rubbery phase containing 14% molar content of PB, indicates a slightly preferential diffusion of cyclopentane in the rubber than in the polystyrene matrix. This is also confirmed by the NMR observation of a significant swelling of the rubber phase, as indicated by relaxation time measurements [85].

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Figure 14.19 Dependence of τ R for the spin probe in gels 1 and 2 on the concentration of CPC. The dashed line represents τ R of free micelles in water. Adapted with permission from [86]. Copyright 2002 American Chemical Society.

The swelling behavior and the structure of composite gels of poly(acrylamide) with incorporated bentonite clay and the products of their reactions with a cationic surfactant, cetylpyridinium chloride (CPC), were studied by the amphiphilic nitroxyl radical V as a spin probe. Single component spectra show that the adsorption of CPC on the clay platelets already leads to the formation of ordered surfactant aggregates with low molecular mobility in the region of surfactant concentrations much below the critical micelle concentration (CMC), and that the rotational mobility of the spin probe solubilized in these aggregates is low in comparison with its mobility in the micelles of the surfactant in the solution [86]. Figure 14.19 shows the dependence of τ R on concentration of CPC for gels 1 (with monomers) and 2 (without monomers). Experimental curves show slight differences in the probe dynamics at low CPC concentration, while at higher concentrations the motional behavior of the probe approaches to that of free micelles in water. A novel method of nitroxyl spin labeling suitable for molecular dynamics studies within the interface regions of styrene-butadiene rubber (SBR)/silica composites by chemical linking of nitroxyl radical III to the surface of SiO2 particles have been developed. The same samples were also studied using nitroxyl radical I as a spin probe. Fast and slow motional components of spin labeled composites simulated by Freed’s program have been identified within the interface regions and the corresponding rotational diffusion tensors have been measured as a function of the temperature and the SiO2 concentration. The fast rotational frequency is found to be orders of magnitude slower than that measured in the absence of SiO2 . This difference is suggested to arise from a closer packing of the macromolecules near the silica surface caused by the van der Waals bonding interactions. Increase of the SiO2 concentration results in a decrease of the molecular mobility. This effect has been imputed to the overlapping of the bonding interaction regions. Spin probe measurements indicate that the hindrance to chain segmental motion induced by the SiO2 -SBR interactions propagates beyond the interface regions thus involving the bulk polymer matrix. It is suggested that the information on the segmental chain dynamics obtained through the spin labeling and spin probe measurements can be developed as a method for investigating the polymer/filler interactions within the reinforcing mechanism by the filler [29]. Spin labeled PEO tethered on silica was studied to characterize the conformation and local dynamics of the chain end. The observed two component spectra arise from two different kinds of spin labels attached to ‘train’ and ‘tail’ segments, which strongly and weakly interact with the silica surface, respectively. The fractional amount of the ‘train’ segments trapped near the silica decreases remarkably with an increase in the grafting ratio [87].

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Polystyrene (PS) composites and polycaprolactone (PCL) nanocomposites containing site-specifically spin labeled surfactants in the organically modified layered silicate magadiite were studied by DEER and ESEEM pulsed ESR techniques. The DEER data prove that even in the case of the non-intercalated PS composites the density of surfactant molecules changes drastically during composite formation on length scales of a few nanometers. Deuterium ESEEM data demonstrate that spin labels attached both in the middle and at the end of the alkyl chain have contact with the headgroups of neighboring surfactant molecules [88].

14.9 Nanocomposites Because of the strong interactions of the matrix chains with filler nanoparticles, nanocomposite polymeric materials exhibit highly improved mechanical, physical and chemical properties compared to the pure polymers. Therefore, they are particularly interesting for the application and fundamental research. ESR spectroscopy is usually used for the study of structure and dynamics of nanocomposite materials [89–91] but also for diffusion measurements [92, 93] and distance measurements in the nanometer range [94]. Natural rubber doped with nanosilica particles was studied by the spin probe method. The crosslinking density was the same in all samples, and the amount of nanofiller was varied. Two-component ESR spectra of nitroxyl radical III measured in a wide temperature range have shown a restricted segmental mobility of the filled samples in comparison with the pristine NR. However, for the higher filler loadings, the nanoparticles agglomerated and consequently no further effect on the restriction in segmental motions was detected. An addition of coupling agent to the samples with lower amount of nanosilica shifted T 5mT value toward higher temperatures. This fact and the result of the analysis of ESR spectra observed in the vicinity of T 5mT suggest a supplementary reduction of chain dynamics due to the addition of a coupling agent. The same experiments ware performed with NR containing organically-modified nanoclay (Cloisite 20A). It was shown that nanoclay has a more pronounced effect on the reduction in segmental motions than nanosilica. Particularly interesting are spectra measured at temperatures above 80◦ C. For such temperatures, three narrow lines indicating fast isotropic motions of spin probe usually occur. However, this is not the case with NR-nanoclay where an unusual line broadening was observed with the temperature increase. The broad lines imply that in such a high temperature region the probe motions might be frozen on the ESR time scale. This effect might be related to the extremely strong interactions of the chain segments with the molten nanoclay [5]. The structure and dynamics of the surfactant layer in polymer–clay nanocomposites were studied using spin labeled surfactants in function of temperature. Three specifically synthesized surfactants were labeled with nitroxyl radicals IV and VII at different positions along their alkyl chains allowing the observation of heterogeneous dynamics and dynamic gradient along surfactant alkyl chain. An addition of polystyrene provoked a decrease in dynamics, and the mobility of the system was dominated by the mobility of the  surrounding polymer matrix. Plotting 2Azz against the temperature made it possible to distinguish between nano- and microcomposites. The results support existing models of the structure, where the surfactants lie flatly on the surface of the clay platelets (Figure 14.20) [91]. Nanocomposites of poly(methyl acrylate) (PMA) with synthetic fluoromica (Somasif) as the inorganic component were studied as a function of clay content by the spin label technique, using nitroxyl radical IV. The properties of these materials were compared with those of the nanocomposites prepared with the clay modified by surfactants. Spectra indicated that the mobility of PMA chains in the nanocomposites is constrained due to the interactions in the interphase region (Figure 14.21). In nanocomposites prepared in the presence of surfactants, the interaction between PMA and Somasif was reduced, and the mobility of PMA was enhanced. No significant effect on polymer dynamics was detected in conventional composites, in which the clay was dispersed on a scale significantly larger than the nanoscale [95].

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Figure 14.20 Structure scheme for organically modified clay. The slow fraction of surfactants is anchored on the clay surface while the more mobile, fast fraction is located between these surface layers. The surfactant tails are almost parallel to the clay surface and exhibit a dynamic gradient along the alkyl chain. Based on [91].

The structure and dynamics of end-labeled PEO intercalated in the galleries of a fluoromica inorganic clay were studied with nitroxyl radical VII. For PEO intercalated in the narrow clay galleries, the spectra indicated a very low segmental mobility even at high temperature, which was attributed to the strong polymer interaction with the charged mica platelets. ESR spectra of all samples were simulated with a single spectral component, with the exception of the sample containing wider galleries. In this sample, the parameters used to simulate the spectra of the nitroxyl labels reflected a lowering of the PEO segmental density and the spectrum consisted of two distinct contributions from slow and fast motional components (the relative intensity of the fast component increased with an increase in temperature). The two spectral components were attributed to segments located close to and away from respectively, the polar solid walls in the gallery. The fast motional component had higher mobility compared to that of PEO chains adsorbed on the fluoromica surface. The low segmental density and reduced cooperative motion with neighboring segments are considered the main factors leading to the fast PEO chain motion with low activation energy [96].

14.10 Other Polymer Multiphase Systems A number of papers related to ESR spectroscopy in multiphase polymer systems could not be included in any of previous categories. Therefore, some of them will be elaborated separately. A part of these studies was focused on liquid crystalline polymers. Andreozzi et al. investigated the heterogeneities induced in a nematic

Nonrestricted Chains

Restricted Chains

Figure 14.21 Schematic presentation of PMMA containing Somasif platelets. ESR spectra show restricted motions in the interphase region. Non-restricted polymer mobility was observed in regions distant from the Somasif surface. Reprinted with permission from [95]. Copyright 2006 American Chemical Society.

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polymethacrylate by thermal annealing in the isotropic phase. They monitored the dynamics of nitroxyl radical XVIII as a spin probe dissolved in the host matrix. The temperature dependences of the dynamics of the slow and fast components were fully characterized in going from the isotropic state to the glassy state through the nematic phase. It was found that the behaviors of the spinning correlation times with temperature in the isotropic and nematic regions were well represented by the Vogel-Fulcher law. The simulations of ESR lineshape combined with a suitable thermal annealing procedure have shown evidence of the presence of different molecular sites in the nematic PMA matrix. The adopted annealing treatment of the nematic PMA at temperatures well above T NI rendered the polymer matrix rich of defects on a nanometer length scale [97]. The same system was also studied in a wide temperature interval after annealing at a temperature value slightly above the T NI [98]. Nitroxyl radicals VI and XVII as spin probes were used to investigate the plasticization of poly(vinyl acetate) (PVAc) and poly(ethyl methacrylate) with carbon dioxide by varying the CO2 pressure at constant ambient temperature. A special high-pressure ESR instrumentation was used to perform the measurements [99]. Extrema separation (2Azz  ) was measured as a function of CO2 pressure. The value P5mT was defined as the CO2 pressure at which 2Azz ’ reaches 5 mT. It was found that the ESR spectroscopy provides higher glasstransition pressure (P5mT ) values because of the frequency of the measurement, compared with traditional methods. The results demonstrated that CO2 has the ability to affect the mobility and free-volume distribution experienced by nitroxyl radical VI at dramatically low concentrations compared with the conventional plasticizer, dibutyl phthalate (DBP) [100]. The heterogeneity of denture resins based on PMMA was studied using nitroxyl radical III as a spin probe. The resins processed by microwave energy, conventional curing and cold curing and exposed to ageing in various environmental conditions were compared. The results based on the free-volume concept indicated that the local density depends on the curing process. The microwave curing contributes to the highest density of the resin with the lowest residual monomer content, while the cold curing produces the least homogeneous resin. The change of the resin structure caused by accelerated ageing was related to both, the original resin structure and the level of residual monomer [101]. ESR spectroscopy was used for the first time to investigate liquid diffusion into contact lenses. The environment of nitroxyl radical I as a spin probe began to change from solid to liquid if liquid were dropped into the system. On the base of ESR results three distinct steps were found and their diffusion coefficients were determined [102]. Nitroxyl radical X was used as spin probe to study the adsorption of a comb-type polymer polymaleic anhydride octyl vinyl ether (PMAOVE) and a diblock copolymer PB-co-PEO onto reverse microemulsion droplets formed from Aerosol-OT/heptane/water. It was concluded that deprotonation/protonation effects of the spin probe could explain the changes in the ESR spectra after addition of PMAOVE polymer, which lower the pH of the reverse microemulsion drop and this leads to protonation of the spin probe. The protonated spin probe is less polar and relocates further away from the water core leading to increased mobility of the nitroxyl radical. Although the presence of PB-co-PEO influences the film rigidity, as found by electrical percolation, these effects cannot be detected by the spin probe method, because of the limited time domain of conventional ESR spectroscpy [103]. ESR spectra of three spin probes with different molecular volumes (nitroxyl radicals II, IX and XIX) in glassy polystyrene, polyvinyl trimethylsilane, and Teflon AF-2400 were calculated numerically within the model of quasi-libration motions. It was established that the average amplitude of quasi-libration motion depends on the free volume of each polymer and geometrical molecular volume of a spin probe. In addition, it was found that upon further temperature increase (above 40◦ C), the quasi-libration model becomes inapplicable for quantitative numerical spectra simulation [104]. The thermoreversible phase transition of poly(N-isopropylacrylamide) (PNIPAM) randomly labeled with nitroxyl radical IV, and a fluorescent dye, 4-(pyren-1-yl)butyl (PNIPAM-Py-T), in different H2 O/MeOH

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mixtures was studied by ESR, fluorescence spectroscopy and turbidimetry. The analysis of the ESR line shapes revealed the different motion characteristics of PNIPAM-Py-T in H2 O/MeOH, and THF. The associated H2 O layer at PNIPAM-Py-T causes the appearance of two ESR components. the analysis of the isotropic hyperfine coupling constant revealed that the hydrophobicity of the precipitate formed from PNIPAM-Py-T is significantly enhanced in comparison to that of PNIPAM-T [105]. A review of spectroscopic probe studies on self-assembling of ionic amphiphiles in aqueous solutions is presented in [106]. Three different types of samples were investigated: (i) ionomers – perfluorinated material which consists of a Teflon-like backbone, and pendant perfluoropolyether chains terminated with sulfonic acid groups, commercially named Nafion, neutralized by lithium hydroxide and protiated ionomer: the copolymer of ethylene and methacrylic acid, EMAA, neutralized by potassium hydroxide; (ii) surfactants – low-molecular weight surfactants include hexadecyltrimethylammonium chloride (HTAC) as a modelprotiated compound, and two fluorinated compounds: ammonium perfluorooctanoate (APFO) and tetraethylammonium perfluorooctanesulfonate (TEAPFOS) and (iii) polyelectrolytes – polymethacrylic acid (PMA) and its block copolymer with dimethylsiloxane (PMA-b-PDMS). The potential of ESR and fluorescence methods is illustrated by the examples of experiments. It has been shown that ESR and fluorescence probe studies provide complementary information on the self assembling of amphiphiles in aqueous solutions, on the internal structure of the aggregates, as well as on the dynamics and location of the probes. An evidence for the existence of equilibrium between large multichain aggregates and small unimeric micelles in aqueous solutions of ionomers has been presented. Micellar dispersions made of poly(ethylene glycol)s–phosphatidylethanolamines (PEGs–PEs) polymerlipids were studied by using nitroxyl radical X as a spin label. The hydrocarbon chain length and the polymer size of the polymer-lipids were varied. Well-defined chain flexibility gradients of the same overall shape are obtained in the considered dispersions. The mobility of the first acyl chain segments was appreciably higher in micelles of polymer-lipids than in bilayers of dipalmitoylphosphatidylcholine (DPPC) and it becomes indistinguishable at the chain termini [107]. The phase-separated morphology and the effect of ionic aggregation on the backbone chain mobility in poly(ethylene-ran-methacrylic acid) (EMAA) ionomers neutralized by Na+ were investigated as a function of the degree of neutralization by spin probe in ESR and by small-angle X-ray scattering (SAXS). Spectra of five spin probes differing in their hydrophilicity and in the position of the nitroxyl group with respect to the headgroup (nitroxyl radicals X, XI, XII, XIV and XVII) were analyzed. The results indicated that the probes are position-selective and can provide information on the local polarity and mobility in, and near, the ionic aggregates. the spectra clearly indicated the highly restricted mobility of the chain segments in the hydrocarbon shell surrounding the ionic aggregates [108]. The poly(acrylic acid) (PAA)/poly(ethylene oxide) (PEO) complex was studied over a wide temperature range (100 K–450 K) with four nitroxyl radicals (I, II, III and IV) as spin probes. The spectra of each probe exhibit two components in certain temperature region. Spectral parameters T 5mT , Ta and Td demonstrate the effects of probe size, polymer matrix rigidity and the strength of hydrogen bonding. T 5mT gradually increases with the probe size due to the free volume decrease, while the molecular motion decreases with increasing content of PAA component [109].

14.11 Conclusion This chapter is an overview of ESR studies of multiphase polymer systems published during the past decade. Earlier investigations in this field are systematically described in [6]. The overview is limited on ESR studies performed by an introduction of nitroxyl radicals into investigated polymer systems. A considerable number of papers representing the ESR studies of biopolymers and homopolymers are not included in this chapter.

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CW ESR spectroscopy in multiphase polymer systems is based on the wide temperature range study of motional and orientational behavior of nitroxyl radicals incorporated in polymer systems. Variable temperature CW ESR measurements usually reveal a bimodal distribution of the rotational correlation times of nitroxyl radicals. This is the consequence of dynamic heterogeneity of multiphase polymer systems. Spectra which consist of two components, broad and narrow, emphasize a coexistence of the slow and fast molecular motion in the system, respectively. Detailed analysis of bimodal ESR spectra made by computer programs described above gives the possibility to distinguish two types of motion and to obtain useful information about local segmental dynamics, spatial orientation, phase behavior and local matrix structure. However, the conventional ESR spectroscopy has a limited time domain and therefore it is not eligible for the study of polymers characterized by slow molecular motions, where no complete time averaging of magnetic interactions occurs. To the contrary, pulsed ESR spectroscopy based on electron spin echo (ESE) is highly sensitive to the changes in molecular motion in the low temperature range (under Tg ). ESE techniques are successful in studying nanostructure matter in the range of ∼3 nm, which cannot be investigated by other methods. The ESEEM through electron–nuclear interactions is a convenient tool of structural studies in the nearest environment of spin labels and therefore it is used for identification of the types of nuclei in the nearest environment and determination of their spatial arrangement relative to the unpaired electron. Recently, this method has found increasing application in studies of the local environment of spin labels. PELDOR, also known as DEER, allows one to determine the distances between radicals on the nanometer scale, get information on the conformations of doubly spin labeled macromolecules and estimate the size of the aggregate of spin labeled molecules. However, the number of published papers dealing with pulsed ESR spectroscopy in multiphase polymer systems is still limited.

References 1. Berliner, L.J., Ed., Spin Labeling II: Theory and Application, Academic Press, New York, 1976. 2. Marsh, D., In: Grell, E., Ed., Membrane Spectroscopy, Springer-Verlag, Berlin, 1981. 3. Weil, J.A. and Bolton, J.R., Electron Paramagnetic Resonance: Elementary Theory and Practical Applications, Wiley-Interscience, Hoboken, New Jersey, 2007. 4. Schlick, S., Ed., Advanced ESR Methods in Polymer Research, Wiley-Interscience, Hoboken, New Jersey, 2006. 5. Vali´c, S., Electron spin resonance in studying nanocomposite rubber materials. In: Thomas, S.E. and Stephen, R.E., Eds., Rubber Nanocomposites: Preparation, Properties and Applications, Wiley-Interscience, Singapore, 2009. 6. Veksli, Z., Andreis, M. and Rakvin, B. ESR spectroscopy for the study of polymer heterogeneity. Progress in Polymer Science, 25 (2000) 949–986. 7. Jeschke, G. Determination of the nanostructure of polymer materials by electron paramagnetic resonance spectroscopy. Macromolecular Rapid Communications, 23 (2002) 227–246. 8. Didovi´c, M., Klepac, D. and Vali´c, S. ESR - spin probe method in studying natural rubber: an effect of the probe incorporation temperature. Macromolecular Symposia, 265 (2008) 144–147. 9. Pace, M.D. and Snow, A.W. Nitroxide spin probe/label study of hydrogen bonding and probe size effects in a linear epoxy polymer. Macromolecules, 28 (1995) 5300–5305. 10. Berliner, L.J., Ed., Spin Labelling: The Next Millennium, Plenum Press, New York, 1998. 11. Bordignon, E. and Steinhoff, H.J., Membrane protein structure and dynamics studied by site-directed spin labeling ESR. In: Hemminga, M.A. and Berliner, L.J., Eds., ESR Spectroscopy in Membrane Biophysics, Springer Science and Business Media, New York, 2007. 12. Kim, S.S., Yavrouian, A.H. and Liang, R.H. Segmental mobility and free-volume decay of poly(phenyl methacrylate). An EPR study. Journal of Polymer Science Part B: Polymer Physics, 31 (1993) 495–499. 13. Marinovi´c, T., Vali´c, S., Andreis, M. and Veksli, Z. Study of matrix inhomogeneity of natural rubber and synthetic polyisoprenes by a spin probe method. Polymer, 32 (1991) 2519–2522.

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14. Vali´c, S., Andreis, M., Veksli, Z. and Charlesby, A. Matrix inhomogeneity in crosslinked rubber and rubber emulsions. Radiation Physics and Chemistry, 37 (1991) 257–261. 15. Kovarskii, A.L., Kulish, E.I., Placek, Y., Szocs, F., Kolesov, S.V. and Minsker, K.S. Application of the spin-probe method for the study of two-phase polymer systems: poly(vinyl chloride)-polyethylene blends. Vysokomolekulyarnye Soedineniya Seriya A, 35 (1993) A1669–A1673. 16. Hlouˇskov´a, Z., Tiˇno, J. and Chod´ak, I. Study of PE-based composites by the spin-probe method. European Polymer Journal, 30 (1994) 175–178. ˇ ˇ 17. Culin, J., Andreis, M., Veksli, Z., Anˇzlovar, A. and Zigon, M. ESR-spin labelling study of semi-interpenetrating networks and polymer mixtures based on functionalized polyurethanes and polymethacrylates. European Polymer Journal, 41 (2005) 1874–1882. 18. Schlick, S., Harvey, R.D., Alonso-Amigo, M.G. and Klempner, D. Study of phase separation in interpenetrating polymer networks using nitroxide spin labels. Macromolecules, 22 (1989) 822–830. 19. Cameron, G.G., Qureshi, M.Y., Stewart, D., Buscall, R. and Nemcek, J. Spin-label study of immiscible polymers: 5. Blends of labelled poly(2-ethyl hexyl methacrylate) with poly(methyl methacrylate) and polystyrene. Polymer, 36 (1995) 3071–3074. 20. Kivelson, D. Theory of ESR linewidths of free radicals. Journal of Chemical Physics, 33 (1960) 1094–1106. 21. Goldman, S.A., Bruno, G.V. and Freed, J.H. Estimating slow-motional rotational correlation times for nitroxides by electron spin resonance. Journal of Physical Chemistry, 76 (1972) 1858–1860. ˇ 22. Dubrovi´c, I., Klepac, D., Vali´c, S. and Zauhar, G. Study of natural rubber crosslinked in the state of uniaxial deformation. Radiation Physics and Chemistry, 77 (2008) 811–817. 23. Kusumoto, N., Sano, S., Zaitsu, N. and Motozato, Y. Molecular motions and segmental size of vulcanized natural and acrylonitrile-butadiene rubbers by the spin-probe method. Polymer, 17 (1976) 448–454. 24. Schneider, D.J. and Freed, J.H., Biological Magnetic Resonance, Plenum Press, New York, 1989. 25. Budil, D.E., Lee, S., Saxena, S. and Freed, J.H. Nonlinear-least-squares analysis of slow-motion EPR spectra in one and two dimensions using a modified levenberg-marquardt algorithm. Journal of Magnetic Resonance, Series A, 120 (1996) 155–189. 26. Stoll, S. and Schweiger, A. EasySpin, a comprehensive software package for spectral simulation and analysis in EPR. Journal of Magnetic Resonance, 178 (2006) 42–55. 27. Stoll, S. and Schweiger, A. EasySpin: simulating cw ESR spectra. Biological Magnetic Resonance, 27 (2007) 299–321. 28. Veksli, Z., Andreis, M. and Campbell, D.S. A spin-probe study of heterogeneity in the natural rubber matrix: The effect of molecular weight, molecular weight distribution and gel phase. Polymer, 39 (1998) 2083–2088. 29. Buttafava, A., Ghisoni, G.M., Faucitano, A., Negroni, G., Priola, A., Peditto, F., Turturro, A. and Castellano, M. EPR spin labelling studies of molecular dynamics in elastomer-silica composites. Research on Chemical Intermediates, 28 (2002) 191–204. 30. Brik, M.E., Titman, J.J., Bayle, J.P. and Judeinstein, P. Mapping of motional heterogeneity in organic-inorganic nanocomposite gels. Journal of Polymer Science Part B: Polymer Physics, 34 (1996) 2533–2542. 31. Veksli, Z., Andreis, M., Vali´c, S., Marinovi´c, T. and Ranogajec, F. Different spatial heterogeneity of networks prepared by a two stage irradiation of natural rubber. Radiation Physics and Chemistry, 51 (1998) 207–213. 32. Freed, J.H. ESR studies of spin probes in anisotropic media. ACS Symposium Series, 34 (1976) 1–15. 33. Schweiger, A. and Jeschke, G., Principles of Pulse Electron Paramagnetic Resonance, Oxford University Press, New York, 2001. 34. Tsvetkov, Y.D., Milov, A.D. and Maryasov, A.G. Pulsed electron-electron double resonance (PELDOR) as EPR spectroscopy in nanometre range. Russian Chemical Reviews, 77 (2008) 487–520. 35. Hamley, I.W. The Physics of Block Copolymers, Oxford University Press, New York, 1999. 36. Shvartzman-Cohen, R., Florent, M., Goldfarb, D., Szleifer, I. and Yerushalmi-Rozen, R. Aggregation and selfassembly of amphiphilic block copolymers in aqueous dispersions of carbon nanotubes. Langmuir, 24 (2008) 4625–4632. 37. Zhou, L. and Schlick, S. Electron spin resonance (ESR) spectra of amphiphilic spin probes in the triblock copolymer EO13 PO30 EO13 (Pluronic L64): hydration, dynamics and order in the polymer aggregates. Polymer, 41 (2000) 4679–4689.

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38. Ruthstein, S., Frydman, V., Kababya, S., Landau, M. and Goldfarb, D. Study of the formation of the mesoporous material SBA-15 by EPR spectroscopy. Journal of Physical Chemistry B, 107 (2003) 1739–1748. 39. Ruthstein, S., Frydman, V. and Goldfarb, D. Study of the initial formation stages of the mesoporous material SBA-15 using spin-labeled block co-polymer templates. Journal of Physical Chemistry B, 108 (2004) 9016–9022. 40. Beghein, N., Rouxhet, L., Dinguizli, M., Brewster, M.E., Ari¨en, A., Pr´eat, V., Habib, J.L. and Gallez, B. Characterization of self-assembling copolymers in aqueous solutions using electron paramagnetic resonance and fluorescence spectroscopy. Journal of Controlled Release, 117 (2007) 196–203. 41. Miwa, Y., Tanida, K., Yamamoto, K., Okamoto, S., Sakaguchi, M., Sakai, M., Makita, S., Sakurai, S. and Shimada, S. Dynamic heterogeneity in interfacial region of microphase-separated polystyrene-block-poly(methyl acrylate) studied by the ESR spin-label technique. Macromolecules, 37 (2004) 3707–3716. 42. Miwa, Y., Yamamoto, K., Tanabe, T., Okamoto, S., Sakaguchi, M., Sakai, M. and Shimada, S. An ESR spin-label study on molecular mobility in the interface between microphases of a diblock copolymer: Effects of admixture of homopolymers that are miscible with one of the blocks. Journal of Physical Chemistry B, 110 (2006) 4073–4082. 43. Andreozzi, L., Faetti, M., Giordano, M., Palazzuoli, D. and Galli, G. Heterogeneity in the dynamics of a molecular tracer dissolved in liquid crystal homopolymer and copolymer. Molecular Crystals and Liquid Crystals, 411 (2004) 515–523. 44. Andreozzi, L., Autiero, C., Faetti, M., Zulli, F. and Giordano, M. Heterogeneities in the dynamics of a molecular tracer in mesogenic and nonmesogenic azobenzene copolymers. Molecular Crystals and Liquid Crystals, 450 (2006) 363–371. 45. Andreozzi, L., Faetti, M., Giordano, M., Palazzuoli, D., Zulli, F. and Galli, G. Effects of thermal annealing on the heterogeneities in the dynamics of a paramagnetic probe dissolved in azobenzene polymethacrylates. Molecular Crystals and Liquid Crystals, 429 (2005) 21–29. 46. Chipara, M., Rowlands, C.C. and Galatanu, A.N. Electron spin resonance investigation of molecular motions in oilextended styrene-butadiene-styrene block copolymers. I. The temperature dependence of resonance spectra: Glass, narrowing, and inflection temperatures. Journal of Polymer Science Part B: Polymer Physics, 42 (2004) 1960– 1971. ˇ ep´anek, P. and Pilaˇr, J. Local segmental dynamics of poly(247. Marek, A., Czernek, J., Steinhart, M., Labsk´y, J., Stˇ hydroxyethyl methacrylate) in methanolic solution by spin label X-band ESR. Journal of Physical Chemistry B, 108 (2004) 9482–9490. 48. Turacchio, M., Di Nino, G., D’Archivio, A.A., Jer´abek, K., Lora, S., Antonini, G. and Corain, B. Nanostructure and molecular accessibility of gel-type resins for supported bio-catalysis. Part I. Poly-hydroxyethylmethacrylatehydroxypropylmethacrylate-trimethylolpropanetrimethacrylate. Reactive and Functional Polymers, 55 (2003) 21–26. 49. Varghese, B. and Schlick, S. Microphase separation in poly(acrylonitrile-butadiene-styrene) (ABS) studied with paramagnetic spin probes. I. Electron spin resonance spectra. Journal of Polymer Science Part B: Polymer Physics, 40 (2002) 415–423. 50. Varghese, B. and Schlick, S. Microphase separation in poly(acrylonitrile-butadiene-styrene) (ABS) studied with paramagnetic spin probes. II. Simulation of electron spin resonance spectra. Journal of Polymer Science Part B: Polymer Physics, 40 (2002) 424–433. 51. Kruczala, K., Varghese, B., Bokria, J.G. and Schlick, S. Thermal aging of heterophasic propylene-ethylene copolymers: Morphological aspects based on ESR, FTIR, and DSC. Macromolecules, 36 (2003) 1899–1908. 52. Kruczala, K., Bokria, J.G. and Schlick, S. Thermal aging of heterophasic propylene-ethylene copolymers: Spatial and temporal aspects of degradation based on ESR, ESR imaging, and FTIR. Macromolecules, 36 (2003) 1909–1919. 53. Motyakin, M.V. and Schlick, S. Spectral profiling by 1D and 2D electron spin resonance imaging: Nitroxide radicals in UV and thermal degradation of poly(acrylonitrile-butadiene-styrene) containing a hindered amine stabilizer. Macromolecules, 34 (2001) 2854–2864. 54. Pilaˇr, J. and Labsk´y, J. Solvent dependence of polystyrene local segmental dynamics in dilute solution by spin-label X-band ESR. Macromolecules, 36 (2003) 913–920. ˇ ˇ ˇ 55. Culin, J., Frka, S., Andreis, M., Smit, I., Veksli, Z., Anˇzlovar, A. and Zigon, M. Motional heterogeneity of segmented polyurethane-polymethacrylate mixtures: An influence of functional groups concentration. Polymer, 43 (2002) 3891–3899.

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ˇ 56. Culin, J., Andreis, M., Veksli, Z. and Gallot, Y. Motional heterogeneity of polystyrene-block-polybutadiene: A spin probe study. Polymer, 44 (2003) 7875–7881. ˇ ˇ ˇ 57. Culin, J., Andreis, M., Smit, I., Veksli, Z., Anˇzlovar, A. and Zigon, M. Motional heterogeneity and phase separation of functionalized polyester polyurethanes. European Polymer Journal, 40 (2004) 1857–1866. 58. Tamano, K., Tanaka, T., Awaga, K., Imae, T., Yusa, S.I. and Shimada, Y. Single-line EPR spectra from radicals encapsulated in aggregates of amphiphilic block copolymers with hydrophobic dendritic pendants in water. Macromolecular Rapid Communications, 27 (2006) 1764–1768. 59. Pannier, M., Sch¨ops, M., Sch¨adler, V., Wiesner, U., Jeschke, G. and Spiess, H.W. Characterization of ionic clusters in different ionically functionalized diblock copolymers by CW EPR and four-pulse double electron-electron resonance. Macromolecules, 34 (2001) 5555–5560. 60. Leporini, D., Sch¨adler, V., Wiesner, U., Spiess, H.W. and Jeschke, G. Confinement effects in ionomers: A high-field pulsed electron spin resonance spectroscopy study. Journal of Non-Crystalline Solids, 307 (2002) 510–516. 61. Virtanen, J. and Tenhu, H. Thermal properties of poly(N-isopropylacrylamide)-g-poly(ethylene oxide) in aqueous solutions: Influence of the number and distribution of the grafts. Macromolecules, 33 (2000) 5970–5975. 62. Virtanen, J., Lemmetyinen, H. and Tenhu, H. Fluorescence and EPR studies on the collapse of poly(N-isopropyl acrylamide)-g-poly(ethylene oxide) in water. Polymer, 42 (2001) 9487–9493. 63. Yamamoto, K., Kyouzuka, S. and Shimada, S. Generation and molecular motion of grafted polymethacrylate chains on an isotactic polypropylene film surface. Macromolecules, 37 (2004) 86–91. 64. Okuda, Y., Hayashi, F., Sakurai, H. and Shiotani, M. Graft polymerization of methacrylic acid onto polytetrafluoroethylene initiated by alkyllithium/electron-donating solvents. Journal of Applied Polymer Science, 94 (2004) 923–931. 65. Qiu, F.R., Chen, S.M., Tan, L. and Ping, Z.H. The study of the miscibility and morphology of poly(styrene-co-4vinylphenol)/poly(propylene carbonate) blends. Polymers for Advanced Technologies, 15 (2004) 453–458. 66. Chen, S.M., Qiu, F.R. and Tan, L. The study of poly(styrene-co-p-(hexafluoro-2-hydroxylisopropyl)-αmethylstyrene)/poly(ethylene oxide) blends. Journal of Applied Polymer Science, 92 (2004) 2312–2317. 67. Qiu, F.R., Chen, S.M. and Ping, Z.H. Study of the miscibility and segmental motion of STMAA-PBMA polymer blends and semi-interpenetrating polymer networks by an ESR spin probe method. Magnetic Resonance in Chemistry, 43 (2005) 411–416. 68. Chen, S.M., Tan, L., Qiu, F.R., Jiang, X.L., Wang, M. and Zhang, H.D. The study of poly(styrene-co-p-(hexafluoro2-hydroxylisopropyl)-α-methyl-styrene)/poly(propylene carbonate) blends by ESR spin probe and Raman. Polymer, 45 (2004) 3045–3053. 69. Miwa, Y., Tanabe, T., Yamamoto, K., Sugino, Y., Sakaguchi, M., Sakai, M. and Shimada, S. Segmental dynamics and self-concentration around chain ends in miscible blend of poly(cyclohexyl methacrylate) and poly(cyclohexyl acrylate) as studied by the spin-label technique. Macromolecules, 37 (2004) 8612–8617. ˇ 70. Culin, J., Gembarovski, D., Andreis, M., Veksli, Z. and Marinovi´c, T. Effect of thermal oxidative ageing on the morphology of natural rubber networks as viewed by ESR. Polymer International, 49 (2000) 845–852. 71. Bolfa, C., Zoleo, A., Sassi, A.S., Maniero, A.L., Pears, D., Jerabek, K. and Corain, B. Cross-linked poly-vinyl polymers versus polyureas as designed supports for catalytically active M0 nanoclusters. Part I. Nanometer scale structure of the polyurea support EnCatTM 40. Journal of Molecular Catalysis A: Chemical, 275 (2007) 233–239. 72. Wasserman, A.M., Kasaikin, V.A., Zakharova, Y.A., Aliev, II, Baranovsky, V.Y., Doseva, V. and Yasina, L.L. Molecular organization and dynamics of micellar phase of polyelectrolyte-surfactant complexes: ESR spin probe study. Spectrochimica Acta Part A: Molecular and Biomolecular Spectroscopy, 58 (2002) 1241–1255. 73. Doseva, V., Yasina, L.L., Aliev, I.I., Vasserman, A.M. and Baranovskii, V.Y. Molecular dynamics and organization of micellar complexes of poly(acrylic acid) with poly(ethylene glycol)-based surfactants. Colloid Journal, 65 (2003) 562–566. 74. Tedeschi, A.M., Busi, E., Paduano, L., Basosi, R. and D’Errico, G. Influence of the headgroup molecular structure on the anionic surfactant-PVP interaction studied by electron paramagnetic resonance of a cationic nitroxide. Physical Chemistry Chemical Physics, 5 (2003) 5077–5083. 75. Hinderberger, D., Jeschke, G. and Spiess, H.W. Network formation involving polyelectrolytes in solution: The role of counterions. Colloid and Polymer Science, 282 (2004) 901–909.

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76. Wasserman, A.M., Yasina, L.L., Aliev, II, Doseva, V. and Baranovsky, V.Y. Molecular structure and dynamics of poly(methacrylic acid) and poly(acrylic acid) complexes with dodecyl-substituted poly(ethylene glycol). Colloid and Polymer Science, 282 (2004) 402–406. 77. Zakharova, J.A., Otdelnova, M.V., Aliev, II, Motyakin, M.V., Wasserman, A.M., Zezin, A.B. and Kabanov, V.A. Effect of organic additives on formation and structure of polyelectrolyte-oppositely charged surfactant complexes. Polymers for Advanced Technologies, 17 (2006) 914–923. 78. Tedeschi, A.M., Busi, E., Basosi, R., Paduano, L. and D’Errico, G. Influence of the alkyl tail length on the anionic surfactant-PVP interaction. Journal of Solution Chemistry, 35 (2006) 951–968. 79. Wasserman, A.M., Yasina, L.L., Motyakin, M.V., Aliev, I.I., Churochkina, N.A., Rogovina, L.Z., Lysenko, E.A. and Baranovsky, V.Y. EPR spin probe study of polymer associative systems. Spectrochimica Acta Part A-Molecular Spectroscopy, 69 (2008) 1344–1353. ˇ ˇ 80. Culin, J., Veksli, Z., Anˇzlovar, A. and Zigon, M. Spin probe study of semi-interpenetrating polymer networks based on polyurethane and polymethacrylate functional prepolymers. Polymer International, 52 (2003) 1346–1350. ˇ ˇ ˇ 81. Culin, J., Smit, I., Andreis, M., Veksli, Z., Anˇzlovar, A. and Zigon, M. Motional heterogeneity and phase separation of semi-interpenetrating networks and mixtures based on functionalised polyurethane and polymethacrylate prepolymers. Polymer, 46 (2005) 89–99. 82. Qiu, F.R., Chen, S.M., Ping, Z.H. and Yin, G.M. ESR spin-label study of poly(styrene-co-methacrylic acid)/poly(εcaprolactone) semi-interpenetrating polymer networks with controlled hydrogen-bond interactions. Magnetic Resonance in Chemistry, 43 (2005) 918–925. 83. Cao, X.D. and Zhang, L.N. Miscibility and properties of polyurethane/benzyl starch semi-interpenetrating polymer networks. Journal of Polymer Science Part B: Polymer Physics, 43 (2005) 603–615. 84. Maddinelli, G., Montanari, L., Ferrando, A. and Maestrini, C. Application of magnetic resonance techniques in investigation of hydrocarbons interaction with composite polymers. Journal of Applied Polymer Science, 102 (2006) 2810–2817. 85. Starodoubtsev, S.G., Ryabova, A.A., Dembo, A.T., Dembo, K.A., Aliev, I.I., Wasserman, A.M. and Khokhlov, A.R. Composite gels of poly(acrylamide) with incorporated bentonite. Interaction with cationic surfactants, ESR and SAXS study. Macromolecules, 35 (2002) 6362–6369. 86. Shimada, S., Maruta, A. and Yamamoto, K. Structure and molecular motion of poly(ethylene oxide) chains tethered on silica, depended on grafting ratio by the spin labeled method. Polymer Journal, 32 (2000) 1038–1043. 87. Mao, Q., Schleidt, S., Zimmermann, H. and Jeschke, G. A pulsed EPR study of surfactant layer structure in composites of a synthetic layered silicate with polystyrene and polycaprolactone. Physical Chemistry Chemical Physics, 10 (2008) 1156–1167. 88. Vandermeulen, G.W.M., Hinderberger, D., Xu, H., Sheiko, S.S., Jeschke, G. and Klok, H.A. Structure and dynamics of self-assembled poly(ethylene glycol) based coiled-coil nano objects. A European Journal of Chemical Physics and Physical Chemistry, 5 (2004) 488–494. 89. Jeschke, G., Panek, G., Schleidt, S. and Jonas, U. Addressing the interface in polymer-clay nanocomposites by electron paramagnetic resonance spectroscopy on surfactant probes. Polymer Engineering & Science, 44 (2004) 1112–1121. 90. Schleidt, S., Spiess, H.W. and Jeschke, G. A site-directed spin-labeling study of surfactants in polymer-clay nanocomposites. Colloid and Polymer Science, 284 (2006) 1211–1219. 91. Cramer, S.E., Jeschke, G. and Spiess, H.W. Measurement of diffusion coefficients of additive molecules in colloidal polymer particles by electron paramagnetic resonance. Colloid and Polymer Science, 280 (2002) 569–573. 92. Ionita, P., Volkov, A., Jeschke, G. and Chechik, V. Lateral diffusion of thiol ligands on the surface of Au nanoparticles: An electron paramagnetic resonance study. Analytical Chemistry, 80 (2008) 95–106. 93. Jeschke, G. Distance measurements in the nanometer range by pulse EPR. A European Journal of Chemical Physics and Physical Chemistry, 3 (2002) 927–932. 94. Panek, G., Schleidt, S., Mao, Q., Wolkenhauer, M., Spiess, H.W. and Jeschke, G. Heterogeneity of the surfactant layer in organically modified silicates and polymer/layered silicate composites. Macromolecules, 39 (2006) 2191–2200. 95. Miwa, Y., Drews, A.R. and Schlick, S. Detection of the direct effect of clay on polymer dynamics: The case of spin-labeled poly(methyl acrylate)/clay nanocomposites studied by ESR, XRD, and DSC. Macromolecules, 39 (2006) 3304–3311.

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96. Miwa, Y., Drews, A.R. and Schlick, S. Unique structure and dynamics of poly(ethylene oxide) in layered silicate nanocomposites: Accelerated segmental mobility revealed by simulating ESR spectra of spin-labels, XRD, FTIR, and DSC. Macromolecules, 41 (2008) 4701–4708. 97. Andreozzi, L., Faetti, M., Giordano, M., Palazzuoli, D. and Galli, G. An ESR study on the heterogeneity of dynamics in a nematic polymer induced by thermal annealing in the isotropic melt. Macromolecules, 34 (2001) 7325–7330. 98. Andreozzi, L., Faetti, M., Giordano, M., Palazzuoli, D., Laus, M. and Galli, G. Fractionary couplings of spin probe to backbone and side group dynamics of a liquid crystal polymer. Macromolecular Chemistry and Physics, 203 (2002) 1636–1642. 99. Batchelor, S.N., Henningsen, B. and Fischer, H. EPR of transient free radicals during photochemical reactions in high temperature and pressure gases. Journal of Physical Chemistry A, 101 (1997) 2969–2972. 100. Harbron, E.J., Bunyard, W.C. and Forbes, M.D.E. Electron paramagnetic resonance spin probe study of carbon dioxide-induced polymer plasticization. Journal of Polymer Science Part B: Polymer Physics, 43 (2005) 2097–2108. 101. Bubalo, V., Jerolimov, V., Bauˇci´c, I., Andreis, M. and Veksli, Z. Influence of accelerated ageing on metha-crylatebased denture resins heterogeneity as viewed by ESR-spin-probe method. Polymer International, 54 (2005) 848–853. 102. Ozgur, E.E., Aydin, E., Ozbey, T., Irkec, M., Bozkurt, B. and Kaptan, H.Y. Investigation of liquid diffusion into contact lenses using an electron spin resonance technique. Journal of Applied Polymer Science, 100 (2006) 2942–2946. 103. Wines, T.H., Somasundaran, P., Turro, N.J., Jockusch, S. and Ottaviani, M.F. Investigation of the mobility of amphiphilic polymer – AOT reverse microemulsion systems using electron spin resonance. Journal of Colloid and Interface Science, 285 (2005) 318–325. 104. Chernova, D.A. and Vorobiev, A.K. Molecular mobility of nitroxide spin probes in glassy polymers. Quasi-libration model. Journal of Polymer Science Part B: Polymer Physics, 47 (2009) 107–120. 105. Ottaviani, M.F., Winnik, F.M., Bossmann, S.H. and Turro, N.J. Phase separation of poly(N-isopropylacrylamide) in mixtures of water and methanol: A spectroscopic study of the phase-transition process with a polymer tagged with a fluorescent dye and a spin label. Helvetica Chimica Acta, 84 (2001) 2476–2492. 106. Szajdzinska-Pietek, E. and Schlick, S. Self-assembling of ion-containing polymers and surfactants in aqueous solutions studied by ESR and fluorescence probes. Journal of Molecular Liquids, 117 (2005) 153–164. 107. Bartucci, R., Belsito, S. and Sportelli, L. Spin-label electron spin resonance studies of micellar dispersions of PEGs-PEs polymer-lipids. Chemistry and Physics of Lipids, 124 (2003) 111–122. 108. Kutsumizu, S., Goto, M., Yano, S. and Schlick, S. Structure and dynamics of ionic aggregates in ethylene ionomers and their effect on polymer dynamics: A study by small-angle X-ray scattering and electron spin resonance spectroscopy. Macromolecules, 35 (2002) 6298–6305. 109. Tan, L., Chen, S.M., Ping, Z.H. and Shen, Y.M. Nitroxide spin probe study of probe size, hydrogen bonding and polymer matrix rigidity effects on poly(acrylic acid)/poly(ethylene oxide) complexes. Magnetic Resonance in Chemistry, 41 (2003) 481–488.

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15 XPS Studies of Multiphase Polymer Systems Mohamed M. Chehimi and Fatma Djouani Interfaces, Traitements, Organisation et Dynamique des Syst`emes (ITODYS), University Paris Diderot & CNRS (UMR 7086), Paris, France

Karim Benzarti Universit´e Paris-Est, IFSTTAR, Paris, France

15.1 Introduction Polymers are used in everyday life in packaging, paints, adhesives, sealants, automotive industry, implants, biomedical diagnostic kits, and so on. Whilst the bulk of these materials offers several advantages such as transparency to light, light weight; resistance to chemical, thermal and mechanical stresses, etc. the surface remains a critical region that must be controlled in terms of physico-chemical properties. For example, for 1 cm3 of cubic polymeric object, and assuming a density of 1 and thus a weight of 1g, the outermost layers have a thickness within the nanometer scale. A thickness of 1–10 nm corresponds to 1–10 mg/m2 ca. 100–1000 ng/cm2 or 0.1–1 ppm for the 1 cm3 polymeric object. Hence the surface is an extremely small portion compared to the bulk. However, in diverse applications, a nanometer scale thickness is more than enough to drastically modify the behavior of materials. For example, surface modification of a polymeric object or a deposition of an ultrathin polymer coating on a given substrate (polymer, metal, metal oxide, semi-conductor) significantly alters the macroscopic behavior of the material under test in terms of adhesion, wetting, antifouling, biocompatibility, corrosion control, to name but a few. As these properties are related to the chemical nature of materials’ surfaces, it follows that it is essential to study polymer surfaces or polymers at surfaces by appropriate sensitive techniques. There is an arsenal of surface sensitive techniques that can be employed to analyze polymers and which contains (but is not restricted to):

r r

electron spectroscopies (AES, XPS, UPS, HREELS) SIMS

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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several reflection IR methods contact angle measurements electron and scanning probe microscopies electrophoretic measurements surface plasmon resonance

The emphasis of this chapter is on XPS, one of the most used electron spectroscopy techniques in the domain of polymeric materials. The success of XPS lies in its possibility to analyze all elements except hydrogen, and the availability of a large amount of data of core level electron binding energies for several elements in diverse chemical environments [1]. Moreover, it is generally non-destructive, quantitative and permits to determine overlayer thicknesses in the nanometer scale. For these reasons, XPS is now routinely used in all aspects of surface science in general and in polymer science and engineering in particular. As this contribution is intended for the nonspecialist who wishes to get acquainted with the technique and its applications to polymers, we shall first briefly discuss its historical background, basic principles and instrumentation. Thereafter, we will review applications to a selected polymeric system (e.g. polymer grafts, polymer blends, interpenetrated polymer networks. . .) taken from the literature and from the authors’ research materials.

15.2 Basic Principles of X-ray Photoelectron Spectroscopy 15.2.1

Photoionization

XPS is based on the photoelectric effect: when we shine monoenergetic soft X-rays (Al Kα or Mg Kα, hν = 1486.6 or 1253.6 eV, respectively) on solid materials (e.g. plates, thin films, powders, fibers, implanted atoms) parked on sample holders in high-vacuum, they emit core and valence electrons as displayed in Figure 15.1. In this photoemission process, and for a light element an electron from the K shell (XPS notation: 1s) is ejected together with electrons from the outer and valence levels. The measured kinetic energy (KE) of the photoelectron is given by: KE = hν−BE−φs −δ

(15.1)

Auger e–

L2,3 or 2p

X-rays photon

(a)

Photoioisation (b)

L2,3 or 2p

L1 or 2s photo e–

L1 or 2s

K or 1s

K or 1s Relaxation (c)

Figure 15.1 (a) Irradiation of a material surface by soft X-rays; (b) induced photoemission process; followed by (c) emission of an Auger electron.

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where hν is the energy of the photon, BE is the binding energy of the photoelectron, φ s is the spectrometer work function and δ is the surface static charge. The latter is always observed with insulating materials such as polymers; it must thus be determined accurately. The binding energy of a core level is characteristic of the emitting element. All elements can be detected by XPS except hydrogen, the 1s electron of which is involved in chemical bonds. Once an element is ionized, the K shell vacancy could be filled by an electron from a higher level with a result of an X-ray emission detectable by either energy or wavelength dispersive spectrometers (EDX and WDX, respectively). The second possibility of relaxation is that the energy corresponding to the L2,3 -K transition is sufficient to eject a second electron from the L2,3 level, called Auger electron, resulting in a doubly ionized state of the element. For light elements, the energy of the Auger electron is approximated by: EKL2,3L2,3 ≈ EK − EL2,3 − EL2,3

(15.2)

The kinetic energy EKL2,3L2,3 of the emitted Auger electron is characteristic of the element irrespective of the beam (X-ray, electron, ion) used to ionize the primary core level. Note that whilst for polymers the Auger peaks are rarely investigated, for metals they are frequently explored for obtaining information on the oxidation state. Indeed, Zn2p3/2 BE is quasi the same for the metal and the oxide while the corresponding Auger lines differ markedly in KE positions. 15.2.2

Surface Specificity of XPS

Because the emitted photoelectrons have low inelastic mean free paths (λ) in the 1–4 nm range (NB: λ is related to KE, see Section 15.2.5 below), it follows that photopeaks arise from the outermost layers only. The analysis depth probed by XPS is given by: d = 3λ cos θ = 2 − 12 nm

(15.3)

where θ is the analysis angle relative to the surface normal. This small depth analysis makes XPS a very surface-sensitive analytical technique. 15.2.3

Spectral Examination and Analysis

A routine XPS analysis usually starts by recording a low-resolution survey spectrum in the 0–1100 eV or more to gain information on the nature of elements present at the surface. It also gives a rough estimation of their relative atomic percents. Survey regions give readily a qualitative picture of any surface chemistry undergone by the material. High-resolution narrow regions are recorded to reveal chemical state differences, for example when a metal or a polymer is oxidized or grafted with a molecular or a macromolecular species that is known to contain the target element but in different chemical environments. We shall consider low and high resolution resolution spectra in the following. 15.2.3.1

Survey Regions: Elemental Analysis

Figure 15.2 displays the survey spectra of poly(ε-caprolactone) film, PCL, and a clean silver sample. Each spectrum is a distribution of photoelectrons versus kinetic energy. However, whilst physicists prefer a spectrum

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3x105

3x106

O1s

Ag3d

PCL

Silver plate

5/2-3/2

C1s 2x106

I (cps)

2x105

O KLL

I (cps)

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3/2

Ag3p

1/2

1x106 Ag4p

1x105

Ag3s

Ag4p Ag4s

O2s

C1s

0

0 0

200

400

600

800

1000

Binding energy (eV)

Figure 15.2

0

200

400

600

800

1000

Binding energy (eV)

Survey spectra of poly(ε-caprolactone) and an argon ion-sputtered silver plate.

plotted against the kinetic energy, surface chemists prefer a plot versus the calculated binding energy as this property is characteristic of the ionized core level in the target element. For PCL surface, the survey spectrum exhibits sharp C1s (285 eV) and O1s (532 eV) peaks, the main O KLL Auger transition peak at ∼ 978 eV (apparent BE corresponding to 509 eV on the KE scale), and also O2s at 27 eV. The binding energy range 0–40 eV corresponds to the XPS valence band. It is extremely useful to distinguish between polymers that have comparable chemical structures such as polyethylene and polypropylene and for which the main C1s regions are quasi identical. The XPS valence bands of PE and PP copolymers and blends were carefully investigated in a quantitative manner in order to estimate the surface fraction of ethylene repeat units [2]. The valence band region can thus be used as a fingerprint of the polymers under investigation and bring valuable structural information despite the very low intensity in this spectral region. Alternatively, Section 15.3.8 gives an example of how to extract from the C KLL Auger region information on sp2 carbon atoms percent in PS-PE block copolymers. In Figure 15.2, the silver spectrum reproduces the electronic structure of the metal since all electrons with energy less than that of the X-ray photon energy are ejected and hence featured in the spectrum. Here, the silver surface exhibits, in the increasing BE order, valence band containing Ag4d, Ag4p, Ag4s, the sharp Ag3d doublet, Ag3p3/2 , Ag3p1/2 and Ag3s peaks. The deeper is the electron level, the higher is the binding energy. Note that the silver sample has adventitious hydrocarbon contamination (C1s) which is unavoidable despite argon ion etching in the UHV chamber prior to analysis. Table 15.1 reports BE values for core level electrons emitted from organic polymers and inorganic materials.

15.2.3.2

Narrow Regions: Chemical Shifts

Back in 1958, the group of Kai Siegbahn in Sweden discovered that Cu2 O (CuI) and CuO (CuII) can be distinguished by XPS (see example in [3]). Indeed, they found a difference in binding energy positions of the respective Cu2p3/2 core electrons which they termed ‘chemical shift’ by analogy with NMR spectroscopy where the chemical shift is usually defined in ppm relative to tetramethylsilane (TMS). For this reason, XPS is also called Electron Spectroscopy for Chemical Analysis (ESCA) and in 1981 Kai Siegbahn was awarded the Nobel Prize of Physics for its development.

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Table 15.1 BE values for core electrons from different elements in polymers and in materials frequently used as substrates of polymer coatings. Core level

Binding energy (eV)

Material

C1s N1s O1s F1s Na1s

285 400 533 690 1072

PE Nylon 6,6 PEG PTFE PSS

Si2p P2p S2p Cl2p Br3d

102 134 168 201 71

PDMS DSPE-PEG PES PVC PBS

C1s Al2p Si2p Fe2p3/2 Au4f7/2

284.6 72.8/74.4 99/103.4 707/711 84

MWCNT Al/Al2 O3 Si/SiO2 Fe/Fe2 O3 Au-coated Si wafer

DSPE-PEG: 1,2-distearoyl-sn-glycero-3-phosphoethanolamine-N-[methoxy (polyethylene glycol)-2000], a PEGylated phospholipid.

The so-called ‘chemical shift’ is the cornerstone of XPS since one can distinguish aliphatic hydrocarbon from carboxylic carbon, sulfides from sulfates, metals from oxides, etc. within a depth of ∼10 nm. Figure 15.3 displays high resolution C1s spectra of PCL, PLGA, PS and a perfluorinated ethylene glycol grafted to gold. The PCL C1s region is fitted with three components, centred at 285, 286.5 and 289.0 eV, assigned to C–C/C–H, C–O and O–C=O (the ester carbon group) chemical environments, respectively. The components are approximately in the 4:1:1 ratio which corresponds to the chemical structure of PCL, hence the quantitative aspect of XPS. In the case of PLGA, a biocompatible and biodegradable copolymer, the C1s peak is fitted with three components that are due to the CH3 group from the lactic acid monomer (C-H at 285 eV), C-O from either glycolic or lactic acid monomers (287 eV) and O-C=O (289.2 eV). The latter corresponds to ester groups from the aliphatic polyester chain that arises from the random copolymerization of glycolic and lactic acids. As far as PS is concerned, the main C1s peak is centred at 284.7 eV. In addition, there is a minor feature at high binding energy side (∼7 eV higher than the reference C-C/C-H carbon type) termed ‘shake-up satellite’. This satellite is observed with aromatic polymers such as polystyrene (PS) and poly(ethylene terephtalate) but also with graphitic carbon. In Figure 15.3, the shake-up satellite is actually located at lower kinetic energy (KE) and accounts for an energy loss that occurs upon the interaction of the outgoing photoelectron with a valence electron; the former shaking the latter up to a higher energy level. In aromatic compounds, the shake-up satellite corresponds to a bonding to anti-bonding transition of the π molecular orbital (π →π * transition) brought by the C1s electrons in the aromatic rings. This satellite can be used as a fingerprint of aromatic compounds. For the perfluorinated graft, the C1s region exhibits a complex structure that highlights a shoulder at 286.5 eV due to C-O bonds from ethylene glycol. It is worth noting the very large chemical shifts due CF2 and CF3 groups and whose C1s components are centred at 291.2 and 293.4 eV, respectively.

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C-C/C-H

4x103

PCL

3x103 I (cps)

2x104 1x104 5x103 0 280

30k

C-O O-C=O

2x103

280

295

285 290 Binding energy (eV)

PS

C-C/C-H

15k 10k

15000 CF2

π – π* 285 290 Binding energy (eV)

295

295

Au-PEGF

20000

20k

280

C-O O-C=O

0 285 290 Binding energy (eV)

25k

5k 0

PLGA

C-C/C-H

1x103

I (cps)

I (cps)

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CF3

285 290 295 Binding energy (eV)

300

Figure 15.3 High-resolution C1s spectra of PCL, PLGA copolymer, PS and a perfluorinated PEG grafted to a gold substrate (Au-PEGF) via ‘click’ chemistry.

Table 15.2 reports a selected list of reference BE values for carbon, nitrogen, oxygen, silicon and sulfur in various organic compounds. As a first approximation, for organics, the binding energy BE is given by: BE = kq + BE◦

(15.4)

where q is the atomic charge, BE◦ the BE for a standard state, here the element with neutral charge, and k is the BE-q relationship slope that depends on the core electron level. The difference BE−BE◦ stands for the chemical shift. q is related to the electronegativity difference of the atoms in the chemical bond under test. As fluorine is more electronegative than oxygen, it yields a higher positive atomic charge for carbon and thus a stronger chemical shift for the corresponding C1s peak (see Figure 15.3). As a rule of thumb, C1s binding energy is shifted by + 1.5 eV per oxygen atom linked to the carbon atom under test. C1s BE in hemiacetal is 288 eV due to two C-O bonds. Similarly, and assuming the effect of C=O on C1s shifts as the effects of two single C-O bonds, the C1s BE is 287.8 eV corresponding to almost twice 1.5 eV. For fluorinated compounds, the C1s BE is roughly shifted by +3 eV per C-F bond. For this reason, C1s BE in CF2 is 291.2 eV (roughly two times 3eV shift relative to the reference value of 285 eV); for CF3 , the effects of three fluorine atoms yield +8.4 eV C1s BE shift relative to the reference value of 285 eV, that 2.8 eV per fluorine atom. From the above, it follows that, in a majority of cases, the chemical shifts induced by neighboring atoms on the target element are additive.

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Table 15.2 Typical binding energies for C1s, N1s, O1s, Si2p and S2p in selected functional groups. Functional group

Chemical structure

Binding energy (eV)

C1s Hydrocarbon Amine Ether, alcohol Hemiacetal (in sugar) Ketone Amide Carboxylic acid, ester Carbamate Carbonate Flurocarbon Difluoro carbon Trifluoro carbon

C-H, C-C C-N C-O-C, C-O-H HO-C-OR C=O N-C=O -COOH, -COOR O-C(=O)-N -O-C(=O)-O C-F F-C-F -CF3

285.0 285.7 286.5 288.0 287.8 288.0 289 289.6 290.5 287.8 292 293–294

N1s Imine (in PPy)a Aromatic N (in P4VP) Amide (in Nylon 6) Carbamate Ammonium Nitro

-N=C -N=C N-C=O N-COO-NH3 + -NO2

398.5 399.3 399.8 400.3 401.5 405.5

O1s Carbonyl Alcohol, ether Ester

C=O, O-C=O C-O-H, C-O-C -O-C=O

532.2 532.8 533.7

Si2p Siloxane (in PDMS) Silicon dioxide

O-Si(CH3 )2 Si-O

101.8 103.4

S2p Sulfide Thioetherb Sulfonatec

Metal-S-(CH2 )n -R C-S −SO3 −

162.5 eV 164 168.2

a deprotonated

pyrrole repeat units in PPy; b in cysteine.c PSS.

BE values for several core levels in all elements (except H) were compiled by Wagner et al. and the data bank is freely available online at http://srdata.nist.gov/xps/ [1]. Beamson and Briggs [4] have published highresolution spectra of 111 homopolymers and reported chemical shifts for C1s, O1s, N1s, F1s, Si2p, S2p, Cl2p and Br3d. Attention was particularly paid to the C1s primary and secondary chemical shifts. The primary chemical shifts are due to the effects of elements directly bound to the carbon atom of interest, while the secondary shifts, the so-called β-shifts, are due to elements in β position. β-shifts are due to strong electron withdrawing functional groups such as ester or carboxylic acid groups. Approximately, for the underlined carbon atoms in the structures C-C-F and C-COO, the chemical shifts relative to 285 eV (C-C/C-H) are ∼1 and 0.5-0.7 eV, respectively.

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Quantification

For a homogeneous solid, the peak intensity (I) for an element A is given by the simplified equation: I ∝ J.T (KE).[A].σ.K.λ

(15.5)

where J is the photon flux, T(KE) is the transmission function of the analyzer, [A] is the concentration of the atom A in the solid, σ is its photoionization cross-section, K is an instrumental factor, and λ is the photoelectron inelastic mean free path. Equation (15.5) can be rewritten: I ∝ s[A]

(15.6)

where s is the sensitivity factor. It follows that the atomic percentages of the elements considered can be given by: [A] at.% = [(IA /sA )]/[(In /sn )] × 100%

(15.7)

where In and sn are the integrated peak areas and the sensitivity factors, respectively. 15.2.5

Determination of Overlayer Thickness

The overlayer thickness is another important property that can be determined from surface analysis in addition to the identification of elements and functional groups. For a substrate covered by a homogeneous and nanometer scale thickness, the Beer-Lambert equation can be manipulated in this regard. If the thickness is too high and far exceeds 10–20 nm, it is necessary to perform a destructive depth profiling using an argon gun. However, this method is not recommended for polymers as it degrades and scrambles their chemical structure and does not merely remove the coating layer by layer. A nondestructive approach, such as angle-resolved XPS, is more appropriate to polymers. For the purpose of this chapter we shall concentrate on the nondestructive method that rests on the Beer-Lambert equation. Let’s start first by examining the change in the spectral shape of a substrate that is covered by a polymer coating. The spectrum displayed in Figure 15.4 shows the spectra of gold plates before and after coating by POEGMA ultrathin layers prepared by surface-confined atom transfer radical polymerization (ATRP). Note that the background has a substantial intensity because it is due to inelastically scattered electrons which appear at apparent high BE side which is actually the low KE side of the spectrum. Indeed, these electrons have suffered KE losses due to inelastic collisions within the solid before reaching the spectrometer. The photopeaks are only due to photoelectrons that have not undergone collisions (outermost layers) or that have had collisions without any energy loss. Compared to the spectrum of clean gold, the background is much more intense. This is because the polymer overlayer drastically attenuates the Au4f doublet whilst the Au4d, 4p and 4s photoelectrons are no longer observed in the spectrum of Au-POEGMA as all of them suffer inelastic collisions and are either absorbed by the polymer overlayer or emitted with energy loss so that they only contribute to the background. Figure 15.4 shows not only a decrease of the Au4f/C1s peak intensity ratio upon attachment of the polymer to the surface but also a significant change in the shape of the background.

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Au4f

1.0x106

Au4d5/2-3/2 O1s 8.0x105

Au4p3/2

I (cps)

C1s

Au4p1/2 Au4s

6.0x105

4.0x105

Clean Au

Au-POEGMA

2.0x105

0.0 0

200

400

600

800

1000

Binding energy (eV) Figure 15.4

Survey spectra of a gold substrate before (clean surface) and after attachment of POEGMA chains.

For the substrate (S) covered by the organic coating (C), the peak intensity for a given core level from the substrate (here gold) is given by: IS = I◦S exp(−d/λ cos θ)

(15.8)

IC = I◦C (1 − exp(−d/λ cos θ))

(15.9)

For the top layer:

where I◦ is the intensity of a core level peak from a clean, homogeneous and infinitely thick substrate, d is the thickness of the polymer coating, θ is the analysis angle relative to the suface normal, λ is the inelastic mean free path defined as the average distance that an electron, with a given kinetic energy, travels between successive inelastic collisions. By considering the susbtrate, the overlayer thickness can be estimated from the peak intensities for a pure and infinitely thick sample, before (I◦ ) and after coating (I): d = λ cos θ. ln(I◦ /I)

(15.10)

The thickness can be estimated in a simple manner provided that the coating is homogeneous in composition and thickness and has a coverage degree = 1. This simple method has the disadvantage of the difficulty to determine I◦ , as the intensity of a given peak depends on several instrumental factors such as the detector efficiency and optimal z position of the sample under the X-ray spot. Also, I◦ corresponds to a clean sample which is also difficult to obtain. Actually, in Figure 15.4 the survey scan of the ‘clean’ gold surface displays a

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very small C1s peak due to adventitious hydrocarbon contamination. Despite thorough in situ argon etching in the spectrometer analysis chamber, this hydrocarbon monolayer builds up automatically as the gold surface is sputtered. Indeed, adsorption of contaminants is driven by surface energy minimization of the substrate which is a high surface energy material as are all metals and metal oxides. Instead of performing the analysis at one single angle, θ can be varied and the peak intensity experimentally determined. For a simple analysis at two different angles θi and θj (with θ < 60◦ ), d can be estimated by: ln(Ij /Ii ) = d/λ.(cos−1 θi − cos−1 θj )

(15.11)

This approach has been used for several years but suffers from the change in the area analyzed as the sample is tilted. In addition, it considers the substrate only, not the overlayer. By considering both the substrate and the coating, IC /IS can be written as a ratio R: R = R◦ [exp(d/λS,C cos θ) − exp(d/cosθ.[1/λS,C −1/λC,C ])]

(15.12)

where λS,C is the inelastic mean free of electrons from substrate when they travel in the coating, and λC,C is that of electrons from the coating and travelling in the top coating itself. Under certain circumstances, one can end up with comparable ineslatic mean free path values, so that λS,C ∼ λC,C = λ and: R = R◦ [exp(d/λ cos θ) − 1]

(15.13)

ln[1 + R/R◦ ] = d/(λ cos θ)

(15.14)

which gives the linear relationship:

where 1/ cos θ is the variable and d the slope to be determined. For metal/oxide systems, one can easily obtain similar λ values as the oxide peak is usually a few eVs close to that of the metal (e.g. Al2p in Al and Al2 O3 ; see (15.5)). For the case shown above (Au-POEGMA), one has to choose a core level electron from gold (e.g. Au4d5/2 ) that has a kinetic energy close to that of the characteristic core electron from the top layer (e.g. C1s). These core level electrons have comparable binding energies and therefore comparable kinetic energies, so that their mean free paths in the organic layer would be similar. A final method to determine the overlayer thickness is to choose two peaks from the substrate that have differing kinetic energy positions. For example, iron emits Fe3p (BE = 56 eV) and Fe2p3/2 (BE = 710 eV) core hole electrons and the difference of 654 eV in KE accounts for a difference in sampling depths. As Fe3p is at high KE position (low BE), its inelastic mean free path and therefore sampling depth (see Eq. (15.3)) is higher than that of Fe2p3/2 which is at low KE position (high BE). It follows that by considering two peaks from a substrate with large difference in BE or KE position, it is possible to deduce the overlayer thickness from the difference in their corresponding λ. values. Using Eq. (15.8) for core level peaks A and B, one can derive the following analytical expression for the overlayer thickness: −1 ◦ ◦ d = sin θ/[λ−1 A − λB ]. ln [(IA /IA )/(IB /IB )]

(15.15)

Equations (15.10, 15.11, 15.14, 15.15) rely anyway on the accurate determination of the inelastic mean free path which primarily depends on KE. However, calculation of the inelastic mean free path is complicated

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and due to matrix effect and elastic scattering, λ is now frequently referred to the ‘attenuation length’. In 1979, Seah and Dench [6] proposed, for substrates coated by organic compounds, that the attenuation length be computed using: λ = (103 /ρ)(49 KE−2 + 0.11 KE0.5 ) in nm

(15.16)

where the kinetic energy KE is in eV and ρ is the density of the organic coating in kg/m3 . However, since Eq. (15.16) was derived from a limited and scattered set of data of λ versus KE, and was found to be inaccurate for high KE values due to elastic scattering (see particularly [7] and [8] ), Cumpson and Seah [9] proposed:  λAL = 0.316a 3/2

 E + 4 nanometres Z 0.45 [ln (E/27) + 3]

(15.17)

where a is a lattice parameter (in nm) that can be estimated from:  a = 10

8

μ ρ NAv

1/3 (15.18)

where Nav = 6.02 × 1023 mol−1 is the Avogadro constant, ρ is the density of the matrix (in kg/m3 ) and μ is the average atomic mass of the matrix (in g) and Z is the atomic number of the matrix. Here, if the top layer is PMMA [(C5 H8 O2 )n ], the average atomic mass number μ would be 6.67 g (averaged over the masses of 5 carbon, 8 hydrogen and 2 oxygen atoms in the MMA repeat units). The values of λ can be computed easily from the website www.lasurface.com, by considering either Seah and Dench or Cumpson and Seah, Eqs ((15.16) and (17.17), respectively). Another approach for calculating overlayer thickness rests on the peak shape analysis method [10, 11] instead of the elastic peak intensity only. In this regard, and over several years, Tougaard has developed a software called QUASESTM (quantitative analysis of surfaces by electron spectroscopy) [12] for the accurate determination of thickness and composition within sample surfaces. The approach of Tougaard is built on the phenomenon that the energy loss structure (observed, e.g. in Figure 15.4 for coated Au) that accompanies an XPS or AES peak carries information on the depth of origin of the detected electrons. The method is nondestructive and therefore allows also studying the change in surface morphology during exposure to various treatments. Tougaard has shown that various structures (layered or pillared materials, island-like deposit, etc.) might give rise to the same peak intensity but not the same peak shape as shown in Figure 15.5 [11]. In the theory of Tougaard, for a measured spectrum J(E), the background correction is given by: Emax F(E) = J (E) − B1 E

J (E  )

E − E dE  [C + (E  − E)2 ]

(15.19)

where F(E) is the flux of electrons excited from atoms with kinetic energy E and taken in the range Ep –  < E < Emax , and C is a parameter. F(E) is the primary excitation spectrum. The factor B1 is adjusted such that F(Ep − ) = 0. E − E accounts for energy loss per unit path length travelled by the photoelectron of interest in the solid. Ep –  corresponds to the low kinetic energy end of the peak and Emax the high kinetic

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b

c

d

25 Cu2p 20

15 d 10 c

5

b a

0 450

500

550

600

Figure 15.5 Four widely different surface structures of copper in gold that give identical peak intensities. Reprinted from [11] Copyright (1998) with permission from John Wiley & Sons.

energy end. J(E) can further be evaluated with the in-depth distribution f(z) by the following equation: 

 J (E) =

dE 0 F(E 0 )

f (z)G(E 0 , z/ cos θ : E)dz

(15.20)

where θ is the emission angle of electrons with respect to the surface normal, and the function G, essentially gives the energy distribution of an electron (with initial energy E0 ) as a function of path length z/ cos θ travelled in the solid. G is of importance in any quantitative analysis of energy spectra of emitted electrons. Equations (15.19) and (15.20) are the basis for quantification by background removal in the spectral regions that extends to ∼ 100 eV below the kinetic energy position of the peak, a region where the background shape highlights lateral and in-depth differences in composition as shown in Figure 15.5. Figure 15.6 illustrates the application of Tougaard’s QUASES to the determination of the thickness of POEGMA graft on gold using the survey scan shown in Figure 15.4 for Au-POEGMA. The inelastic background in the O1s region of the survey scan is fitted with QUASES. The darker area in the box indicates the active part of the specimen, here the top organic layer POEGMA. It is assumed that oxygen is mainly due to POEGMA. The dark box indicates the start and end of depth at which POEGMA lies; here between 0.5 and 15.4 nm. The top 0.5 nm correspond probably to a surface contamination layer. The best fit was ˚ using the method of Tanuma et al. [13] and assuming the obtained with O1s electron mean free path of 26.5 A

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0

I (cps)

Depth (nm)

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810 810

850 850

890 890

Kinetic energy (eV)

930 930

970 970

Figure 15.6 QUASES analysis of the O1s region from Au-POEGMA which shows the POEGMA layer is about 15 nm thick.

polymer cross section [14]. The detection of gold Au4f peaks despite a thickness of 15.4 nm is probably due to scratches in the film which are visible by AFM (figure not shown). Very recently, Gam-Derouich et al. [15] have shown that for PHEMA grafts on gold substrates, peak fitting the background with QUASES permitted to look to the very well buried surfaces. The value of 100 nm found for the organic coating (upper limit of thickness obtainable via QUASES) was consistent with the important heights (up to 300 nm) observed in the AFM of PHEMA-coated gold plates. The QUASES method has been applied particularly to inorganic surfaces to study growth, structure and morphology of surface oxides [16], alkali halide model insulator thin layers [17], iron/semiconductor bilayers [18], and supported nanoparticles [19], among other numerous examples. However, the application to polymer surfaces is still sparse, and concerned the study of, for example, plasma polymer surface modification [20], Au nanocluster growth mechanism on polystyrene [21], polymer grafts prepared by SIPP [15]. Recently, L´opezSantos et al. [22] summarized case studies of inorganic and polymer surfaces using QUASES. It is thus clear that given its implementation within the new generation XPS data processing softwares, its commercial availability and the possibility of using it in a friendly manner by XPS users, applications of Tougaard’s QUASES will expand in the next years. Particularly, it has been recently extended to the construction of 3D images of surfaces at the nanoscale [23]. 15.2.6 15.2.6.1

Instrumentation Spectrometer Design

A schematic representation of a modern XPS machine with its primary components is shown in Figure 15.7. To run XPS measurements on polymers and other materials, most modern machines have the following components:

r

An ultrahigh vacuum (UHV) chamber pumped to pressure as low as 10–10 –10–8 mbar using a Turbo or ion pump. A fast entry-lock is also needed to pump, store and pretreat samples. Its pressure is lower than 10–6 mbar.

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Electron energy analyser

Detector X-ray source

Sample Introduction Preparation chamber

Sample

UHV chamber Computer for data acquisition

Vacuum Vacuum Pump Pump

UHV Pumps

Figure 15.7

r r r r r r

r

Schematic diagram for an XP spectrometer showing the required basic components.

X-ray source (polychromatic AlKα/MgKα twin anode, or monochromated AlKα). Electron analyzer which can disperse the emitted electrons according to their kinetic energy. Lens system to ‘drive’ the emitted photoelectrons from the sample to the analyzer. Multi-channel detector. Multi-stage XYZ manipulator to place the samples under the X-ray spot. Electron flood-gun to compensate for static charge built up by insulators. This item is mandatory for the analysis of conventional (but not conductive) polymers, particularly when a monochromated X-ray source is used. Note that the electron flood gun can be used in situ with a stencil mask in order to produce crosslinked monolayers [24] or structured polymer brushes [25]. The technique is called electron beam chemical lithography (EBCL). Computer with XPS software for data acquisition and processing. All XPS machines are now computercontrolled.

Note that on state-of-the-art XPS machines, a parallel imaging facility is available, enabling chemically resolved images with a lateral resolution of 1 μm or less to be obtained. Parallel imaging is nondestructive, facile and very fast. However, there are several methods of obtaining 3D images. Their principles, advantages and limitations were recently reviewed [26]. Whenever destructive depth profiling is needed to analyze the buried layers, an argon ion gun must be installed on the UHV analysis chamber.

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599

Data Acquisition and Processing

There are standard conditions for recording spectra. For survey regions (low resolution spectra) the spectra are recorded with a step size of 1 eV, pass energy of 100–150 eV, in the 0–1000/1100 eV, and dwell time 100 ms. Usually survey spectra are scanned 3–4 times in order to improve the signal-to-noise ratio. For the narrow regions, high resolution spectra can be recorded with a pass energy set at 10–40 eV, a step size of 0.05–0.2 eV, dwell time 100 ms. Spectra are scanned several times, up to 100 scans if a high signal-to-noise ratio is needed at the ultimate resolution or if the surface content of a given element is very low (below 1%). Data processing consists of the following points (all or part of them):

r r r

spectral calibration, e.g. using the so-called internal standard consists in setting a known core level peak at standard position (Au4f7/2 from gold at 84.0 eV or C-C/C-H C1s component from a polymer set at 285.0 eV) background removal using a linear, Shirley ‘s-shaped’ or Tougaard background quantification that gives an apparent composition by considering the peak area and the corresponding Scofield photoionization cross sections corrected for the analyzer transmission function. In the actual commercial software (e.g. Avantage from Thermo), quantification rests on the normalization of peak areas to the sensitivity factor taken as: s = σ.T (KE).λ ≡ σ.T (KE).KE0.5

(15.21)

assuming that λ is a function of KE0.5 . It follows that for elements A and B, their calculated atomic ratio: 0.5 A/B = [IA /σA .T(KEA ).KE0.5] A ]/[IB /σB .T(KEB ).KEB

r

(15.22)

Spectra can be fitted by the least squares method with as many components as necessary. All components must be defined at least by four parameters: peak position, peak height, full width at half maximum, and a Lorentzian/Gaussian shape. For polymers, usually the components have Gaussian shapes, the Lorentzian proportion is small (0–20%). Examples of fitted spectra with minimal χ 2 are shown in Figure 15.3 for PCL and PLGA. Note that the best fit is obtained for the minimal χ 2 but this must make spectroscopic and chemical sense. The best fit must thus reflect a physicochemical reality, not merely a mathematical solution. For further information, the reader is referred to Ref. [7], Appendix 3.

15.3 Applications of XPS to Polymeric Materials Although in the review of the basic principles of XPS we tackled some aspects of XPS analysis of polymers (elemental analysis, chemical shifts and overlayer thickness), we shall now present selected XPS results that are of importance for the understanding of macroscopic, physicochemical behavior of polymers. These applied aspects are restricted to:

r r r

vinylic polymer grafts colloidal particles epoxy adhesives

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conductive polymer materials polymer blends nanocomposites interpenetrated polymer networks random and block copolymers.

There is also abundant literature on XPS studies on plasma deposition of polymers and plasma treatment of polymers. However, these aspects will not be covered here and the reader is referred to [27] and cited literature therein. 15.3.1

Polymer Grafts

It is well known that polymer grafts impart a wide panel of properties to surfaces, including adhesion, wetting, lubrication, response to external stimuli, corrosion control, biocompatibility, bioactivity and antifouling properties to name but a few. Although polymer coatings deposited by solvent casting or by spin coating might have sufficient adhesion to substrates, lifetime durability could be short and adhesion may fail. Grafting polymers to surfaces is thus important to achieve. This can be done either by casting polymers with reactive chain ends or by surface-confined polymerization as schematically shown in Scheme 15.1. Recent years have witnessed a quantum jump in surface-confined controlled radical polymerization (CRP) methods [28, 29] including atom transfer radical polymerization (ATRP), reversible addition-fragmentation transfer polymerization (RAFT) and nitroxide mediated polymerization (NMP) as they permit to grow polymer chains with narrow size distribution and controlled thickness. In order to confine these polymerization techniques to surfaces, it is necessary to pretreat the latter by polymerization initiator-containing surface modifiers (e.g. silanes, thiols and aryl layers). As the thickness remains in the nanometer scale but sometimes up to the thin film regime (submicrometer thick), XPS is an adequate tool to monitor polymer growth on surfaces. Concerning specifically ATRP, XPS has been used to: (i) testify for the attachment of initiators [30, 31]; (ii) characterize the chemical structure of polymer grafts [32, 33], and the living character of the polymer grafts [34]; (iii) estimate the average molecular weight of grafts and monitor their controlled growth. Applications of polymer grafts include pH- [35] and thermal-responsive [36] materials, adhesion properties [37], antifouling properties [38], controlled bacterial adhesion [39], etc. In developing a new approach for production of polymer grafts conductive surfaces, we at University Paris Diderot investigated the possibility of initiating ATRP by brominated aryl layers attached to the substrates by electrochemical reduction of the parent diazonium salts. Several polymer grafts were prepared on iron [30, 33, 34], gold [40], sp2 carbon plates [32, 33] and fibers [41], ultrananocrystalline diamond [42], etc. Concerning reactive grafts, Nguyen et al. [33] prepared PtBMA grafts on glassy carbon (GC) plates via tandem chemistry of diazonium salts and ATRP. The PtBMA grafts were further hydrolyzed in order to convert the ester into carboxylic acid groups (PMAA). GC-PMAA hybrids were found significantly more hydrophilic than GC-PtBMA precursors, with an average water contact angle of 42.3◦ , lower than 87.2◦ found for the latter. XPS was used to monitor changes in the surface composition of GC by PtBMA and to track silver ion retention by the pristine PtBMA grafts as well as by their hydrolyzed form PMAA (Figure 15.8). The polymer grafts impart a sharp O1s peak due to the ester group in the tBMA units (Figure 15.8(a)). The carboxylic acid-functionalized GC-PMAA hybrids were evaluated in the uptake of silver ions from silver acetate aqueous solutions. Figure 15.8b shows that GC-PMAA hybrid is much more effective in removing silver cations than GC-PtBMA as qualitatively judged from the Ag3d/C1s intensity ratios. The selective uptake of Ag+ was interpreted in terms electrostatic interactions between the partially negatively charged carboxylic acid groups (in the form of –COO– ) and the silver cations.

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grafting onto

Adsorbed Polymer

adhesive endgroup

antifouling polymer

601

initiating group

initiator Immobilization

Surface-Initiated Polymerization (SIP)

grafting from

Polymer Brushes

Scheme 15.1 Methods for preparing ultrathin polymer coatings. ‘Grafting onto’ consists of attaching a polymer with a reactive end to the substrate, while ‘grafting from’ consists of functionalizing the surface with a polymerization initiator (a bifunctional molecule or macromolecule) before proceeding to the in situ polymerization for the growth of a thin polymer film.

In a similar way, Mahouche Chergui et al. [43] prepared chelatant polymer grafts on carbon fibers via ATRP surface-initiated by aryl layers derived from diazonium salts. The resulting adsorbents consisted of carbon fiber-grafted cyclam-functionalized poly(glycidyl methacrylate) chains (CF-PGMA-Cy). The adsorption maximum capacity of the CF-PGMA-Cy fibers for Cu(II) was found to be 28.6 mg/g at pH 5.2 as judged from adsorption isotherms determined by the depletion method in conjunction with inductively coupled plasma (ICP) analysis of copper in the supernatants. The optimal pH for adsorption determined by XPS analysis of CF-PMGA-Cy after uptake of copper was also found to be the same, i.e. 5.2. Figure 15.9 shows XPS-determined Cu/N ratio versus pH with an optimum at pH 5.2. Here, N was used as a unique elemental marker for the chelatant cyclam moiety borne by PGMA grafts. The XPS survey scan (insert) highlights Cu2p doublet (in 930–960 eV region) retained by CF-PGMA-Cy fibers. Of relevance to cardiology, Shaulov et al. [44] grafted PMMA chains to the surface of stainless steel (SS) stents via ATRP initiated by diazonium salts. XPS brought supporting evidence for aryl layer grafting followed by attachment of PMMA. For the stent-grafted PMMA, the surface C % was found to be 76.1%,

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C1s

Intensity (a. u.)

Intensity (a. u.)

O1s

GC-PtBMA 5h

Ar2p

0

200

Ag3d GC-PMAA

GC-PtBMA

ion etched GC

400

600

800

0

1000

200

Binding Energy (eV)

400

600

800

1000

Binding energy (eV)

Figure 15.8 Survey spectra from ion-etched GC plate and GC-PtBMA prepared after 5 h ATRP (top). Note that the GC plate was cleaned by argon ion etching. Implanted argon gave an Ar2p peak at ∼241 eV. After hydrolysis of PtBMA grafts into PMAA grafts, uptake of silver cations by GC-PMAA was found more efficient than by GC-PtBMA hybrids (bottom). (Top figure was adapted from Ref. [33]; bottom figure was reproduced from Ref. [33]). Reprinted from [33]. Copyright (2008) with permission from John Wiley & Sons.

much higher than that found for the stent coated from a PMMA solution (60.5%). In addition, the C 1s peak component appearing at 289.3 eV was assigned to the ester bond in the PMMA-grafted layer. It contributed 22.8% to the total carbon C1s peak, very close to the theoretical value of 20% for pure PMMA. XPS was thus very efficient in probing the important chemical changes at the surface of stents and permitted to understand the biomedical applications of these implants. Indeed, the stent-grafted PMMA

C1s

0.4

I (a.u.)

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0.3

O1s

Cu2p

N1s pH=5.2

0.2 0

400

800

1200

XPS-determined copper

Binding energy (eV)

0.1

0.0 1

2

3

4

5

6

pH

Figure 15.9 Uptake of copper ions by chelatant CF-PGMA-Cy fibers versus pH. XPS-determined Cu/N atomic ratio versus pH effect on copper uptake by CF-PGMA-Cy as determined by XPS. Reprinted from [33]. Copyright (2008) with permission from John Wiley & Sons.

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was spray-coated with a drug-in-polymer matrix poly(n-butyl methacrylate)/poly(ethylene-co-vinyl acetate) containing paclitaxel, an antitumoral drug. This approach based on covalently bonded PMMA to the stent, via the aryl layer, was found to significantly improve the durability of drug-in-polymer matrixes on SS drug eluting stents. Instead of CRP methods, much attention has been paid recently to surface-initiated photopolymerization (SIPP) [15, 25, 45] because it is much faster and can be applied to several curved, lamellar and planar substrates. In this regard, Gam-Derouich et al. [15] proposed a novel approach for reactive and functional polymer grafts based on the surface modification of metals by electroreduction of diazonium salt of the 4-benzoyl phenyl group (–C6 H4 -C(O)-C6 H5 , BP in short). BP-modified gold plates served for the SIPP of styrene, methyl methacrylate and 2-hydroxyethyl methacrylate. The resulting PS, and PMMA and PHEMA were characterized in terms of composition and thickness (determined via QUASES) by XPS. They were further subjected to nonspecific protein adsorption of bovine serum albumin (BSA), a model protein. Examination of the high-resolution C1s regions highlighted features from BSA. Indeed, the spectra displayed in Figure 15.10 were peak-fitted with 4 components centred at 285, 286.5, 288 and 289 eV corresponding to the chemical environments C-C/C-H, C-O/C-N, C=O/N-C=O, and O-C=O, respectively. For PS, there is a π –π * satellite at 291.5 eV. For BSA, the most important components are the C–O/C–N and the N–C=O components at 286.5 and 288 eV, respectively. These peak components are noted C286.5 and C288 . In comparison to the C1s spectrum shown in Figure 15.3 for PS, clearly massive spectral changes hold for the substrate Au-BP-PSBSA; it displays prominent C286.5 and C288 components assigned to the adsorbed BSA. The C288 feature due to N–C=O amide groups in proteins is also visible in the spectrum of Au-BP-PMMA-BSA indicating

2500

16000 Au-BP-PS-BSA

BSA

14000 C-N/C-O

1500

I (cps)

I (cps)

2000 N-C=O 1000 500 0

280

16000

285 290 Binding energy (eV)

C-N/C-O

10000

N-C=O

6000 280

295

Au-BP-PMMA-BSA

12000

285 290 Binding energy (eV)

295

Au-BP-PHEMA-BSA

10000

14000 I (cps)

12000

8000

12000

C-N/C-O

10000

N-C=O

I (cps)

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8000 280

2000 285 290 295 Binding energy (eV)

280

285 290 Binding energy (eV)

295

Figure 15.10 Peak-fitted C1s regions from BSA; and Au-BP-PS, Au-BP-PMMA and Au-BP-PHEMA after BSA adsorption. Reproduced from Ref. [11].

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adsorption of BSA onto PMMA grafts. In contrast, the spectrum of Au-BP-PHEMA-BSA is fitted with three main components all due to functional groups from PHEMA (C-C/C-H, C-O and O-C=O at 285, 285.6 and 289 eV, respectively) without any sign for the C288 component characteristic of proteins. This is a strong supporting evidence for the antifouling behavior of PHEMA grafts as processed by the new protocol of SIPP using diazonium salt initiators. The elemental and chemical changes of the polymer surfaces were related to their hydrophilic/hydrophobic nature. BSA adsorption decreased in the order PS > PMMA > PHEMA which is the decreasing order of hydrophobic character as measured by contact angles of water drops deposited on pristine polymer grafts. Finally, concerning polymer grafts it is interesting to report on the possibility of preparing patterned brushes. To do so, one can remove selectively the initiator groups from the surface either electrochemically [46], or using a flood gun in situ in the XPS as mentioned above [25]. In their study of patterned polymer brushes, Slim et al. [46a]; have used imaging XPS to check for the effective elimination of grafted initiator resulting in localized surface-initiated ATRP and therefore formation of patterned brushes. The samples were also characterized by a line-scan procedure permitting to give a surface chemical composition over a patterned region. Figure 15.11 illustrates XPS analysis in the spectroscopic or the imaging modes in the case of a silicon surface grafted with poly(2-(dimethylamino)

(a)

(c)

Ols a

c, d

CIs Si2s Si2p3/2

Counts

NIs

10 b

0

500 400 300 200 100

XPS Image N1s

0

Binding Energy (eV) (d)

(b)

350 3000

250

2500

200 150 100

1600 1200

2000 1500

800

1000 400

500

50 0

0

0

75 150 225 300 375 450 X (μm)

0

N1s (counts)

300 Si2p (counts)

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Figure 15.11 XPS analysis of a silicon plate grafted with patterned PDMAEMA grafts. (a) Illustration of the XPS analyses: gray zones are the patterns and the dashed rectangles the domain analyzed by XPS. (b) XPS image of the N1s photoelectron of a 150 μm width oblique pattern. Black parts were locally reduced to remove the initiator and are thus deficient in N. (c) Survey spectra from a pattern line scan (step size 50 μm). (d) Si2p ( ) and N1s () peak intensity versus position of the silicon plate exposed to the X-ray spot. Reproduced from Ref. [46a].

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605

ethylmethacrylate), PDMAEMA, a pH-responsive and bactericide polymer. In the imaging mode, the parallel N1s image (Figure 15.11(b)) of a straight pattern (ca. 200 μm wide) shows that ATRP confined to the locally etched surfaces yields much lower N surface concentration than the region that has not been reduced. In the spectroscopic mode, Figure 15.11(c) shows the lateral variation of XPS spectra taken over a few millimetres-line perpendicularly to a large patterned area. The N1s peak intensity decreases while the Si2p peak intensity increases when going from the outside to the inside of the pattern (Figure 15.11(d)). These results are interpreted owing to the small mean free path of photoelectrons in solids (generally <2-4 nm). If the layer of the ATRP initiator (grafted from a silane) is not reduced, a thick polymer brush (26 nm) is built that screens the XPS Si2p signal of the silicon substrate. If the bromo-terminated organic layer is reduced, the lower N1s signal indicates that polymerization occurs in the pattern, but results in a thinner and less densely grafted polymer layer (<10 nm).

15.3.2

Colloidal Particles

Analysis of the surface composition of colloidal particles has long attracted colloid scientists and industrials. In this regard, XPS has been used to understand the charge or steric stabilization of colloidal particles by ionic groups or polymeric stabilizers, respectively (e.g. sulfates and PNVP). Several research groups have contributed to the study of surface composition and reactivity of colloidal particles using XPS and other techniques (including SIMS analysis) [47–51]. Particularly, colloidal particles designed for biomedical diagnostic assays have usually surface reactive groups for linkage of biomacromolecules [52]. These specific binding sites can be brought by a comonomer introduced in the synthesis stage or by post-functionalization of prepared particles. However, care must be taken in order to bind specific biomacromolecules and limit their non-specific adsorption (e.g. via hydrophobic interactions). Towards this end, Basinska et al. [53] designed core/shell colloidal particles of poly(styrene-co-αtertbutoxy-ω-vinylbenzyl-polyglycidol) microspheres, hereafter denoted P(S/VB-PGL), by emulsion copolymerization of styrene and VB-PGL, the latter being a hydrophilic macromonomer (see Scheme 15.2 for preparation). The size, bulk composition and surface charge/composition of P(S/VB–PGL) depend on the initial concentration and molecular weight of the hydrophilic VB-PGL macromonomer. It was found that increasing the molar feed ratio of VB-PGL macromonomer (either VB-PGL with MW = 950 or 2700) resulted in smaller microspheres because the macromonomer units could act as stabilizers for the growth of the microspheres. Detection of these hydrophilic units at the surface is thus important and XPS was useful for understanding resistance of P(S/VB-PGL) to nonspecific protein adsorption. Figure 15.12 depicts survey scans and high resolution C1s regions of P(S/VB-PGL) and PS. The main peaks in Figure 15.12a correspond to C1s (285 eV) and O1s (532 eV). Additionally, there is a minor S2p peak centred at ∼169 eV which is assigned to the persulfate initiator of the styrene and VB-PGL copolymerization. O1s peak is due to the sulfate initiating sites in the case of pure PS particles whilst for P(S/VB-PGL) it has a much higher relative intensity due to the polyglycidol-rich surface of P(S/VBPGL). Figure 15.12b shows the high resolution C1s regions from the P(S/VB-PGL) and the reference PS, both exhibiting a main component at 285 eV and a minor one at 291.6 eV assigned to the π →π * shakeup satellite, a ‘fingerprint’ of the aromatic rings in styrene and VB-PGL repeat units. The main difference between the C1s regions lies in the C-O-C C1s component at 286.7 eV resulting from the polyglycidol chains of P(S/VB-PGL). The PGL surface fraction of the microspheres was determined by considering the total C1s and O1s peak areas, the contribution of C-O carbon atoms from the C1s peak fitting (component centred at 286.7 eV), and

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Handbook of Multiphase Polymer Systems CH

CH2

CH

CH2

CH2

(OCHCH2)n~35 OC(CH3)3

1. oxalic acid 2. Ca(OH)2

CH2

(OCHCH2)n

OC(CH3)3

CH2 O H3C

CH2OH PGL-ether

PGL

CH

styrene K2S2O8

O

H2O

CH2CH3

emulsion polymerization

PGL-rich surface

Core/shell P(S/PGL) particles

Scheme 15.2 Synthetic route for core/shell P(S/PGL) microspheres.

C1s

(b) (a) C-C/C-H P(S/PGL) - - - PS

I (a.u.)

O1s

S2p

P(S/VB-PGL)

I (a.u.)

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C-O π-π*

PS

0

200

Figure 15.12

400 600 Binding energy (eV)

800

1000

280

285 290 Binding energy (eV)

295

Typical survey scans (a) and high resolution C1s regions (b) taken from P(S/VB-PGL) and PS.

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XPS Studies of Multiphase Polymer Systems

607

0.5

PGL surface fraction

0.4

0.3

0.2

0.1

0.0

0

1

2 3 4 Bulk PGL fraction (mmol/mol)

P(S/VB-PGL950)

5

6

P(S/VB-PGL2700)

Figure 15.13 Fraction of PGL units in the surface layer of the P(S/VB-PGL) microspheres versus the concentration of PGL in bulk (in mmol/mol) for P(S/VB-PGL950) and P(S/VB-PGL2700) microspheres (plotted using numerical data reported in Ref. [53]).

the fractions x, y and z shown in Figure 15.13 corresponding to styrene, VB-PGL, and PGL units in the macromonomer, respectively. The expression for the C1s peak area is: IC1s α sC [8x + (13 + 3z)y] and since the fraction y = 1 − x, then IC1s α sC [8x + (13 + 3z).(1 − x)] Indeed in the VB-PGL macromonomer there are 13 carbon atoms and z times the three carbon atoms in the glycidol short chain -(O-CH(CH2 OH)-CH2 )z -. Because the oxygen content brought by the sulfate is negligible, one can write for the O1s peak area: IO1s α sO (1 − x).(1 + 2z) Each VB-PGL repeat unit has two oxygen atoms per PGL unit plus one terminal t-butoxy oxygen atom (-O-C(CH3 )3 ). From the experimental C1s and O1s peak areas and the corresponding sensitivity factors, the fraction y of the VB-PGL macromonomer is given by:

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Handbook of Multiphase Polymer Systems 8 y=1–x=

(CH

αO ∑S(C) (1 + 2z) – 3z – 5 αC ∑S(O)

CH 2 ) x

(CH CH 2 ) y

CH 2 (OCHCH 2 ) x OC(CH 3 )y CH 2 OH

The fraction of PGL units, f(PGL), can further be determined using: f(PGL) = yz/[y(z − 1) + 1] The plot shows that for both types of microspheres the surface fraction of PGL units far exceeds that in the bulk which takes into account all styrene and VB-PGL repeat units. The microspheres with PGL-rich surface were evaluated in terms of resistance to nonspecific protein adsorption. Figure 15.14 depicts a plot of the maximum protein adsorption ( ads ) determined for each microsphere versus the corresponding surface PGL fraction as determined by XPS. Clearly, regardless of the molecular weight of the macromonomer, both P(S/VB-PGL) hydrophilic particles exhibited a better resistance to nonspecific human serum albumin (model protein) adsorption compared to the hydrophobic PS particles.

1.6 Maximum HSA adsorption (mg/m2)

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1.2

0.8

0.4

0.0 0.0

0.1

0.2

0.3

0.4

0.5

VB-PGL surface fraction (XPS) PS

P(S/VB-PGL950)

P(S/VB-PGL2700)

Figure 15.14  ads determined for each microsphere versus XPS-determined surface PGL fraction (based on numerical data reported in Ref. [53]).

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XPS Studies of Multiphase Polymer Systems CH2

609

CH O

Scheme 15.3 Structure of oxirane group.

15.3.3

Epoxy Adhesives

Epoxy resins constitute a broad family of compounds containing the oxirane ring in their molecular structure (Scheme 15.3). This functional group reacts readily with compounds possessing active hydrogen atoms, including aliphatic primary amines, tertiary amines, aromatic amines, amides, acids and anhydrides. The choice of crosslinker depends on the application requirements. Epoxy resins have been commercially available for 60 years; the majority of them are based on the diglycidyl ether of bisphenol A (DGEBA) which can be synthesized as shown in Scheme 15.4. Epoxy resins have exceptional properties, such as good adhesion to many substrates; moisture, solvent and chemical resistance; low shrinkage on cure; outstanding mechanical and electronic resistance properties. For this reason they have several industrial applications (as adhesives, surface coatings, laminates and matrix materials in electronic, transport and aerospace industries, civil engineering) but higher costs compared to other thermosets. Typical reactions between epoxy and amine hardener are shown in Scheme 15.5; the kinetics of cure varies with the ratio of hardener to epoxy resin. The successive reactions give rise to a large variety of functional groups that can be probed by XPS. For this reason, XPS has routinely been used to investigate the interfacial chemistry of epoxy resins: metal/epoxy interactions including covalent bonding, the influence of the metallic substrate nature on the prepolymer crosslinking, catalytic effect of the substrate on the crosslinking mechanisms and the prepolymer adsorption onto metallic surfaces [55]. Other studies have attempted to determine the existence of monomer/substratespecific reactions. For example, Dillingham and Boerio [56] have studied the polymerization of DGEBA and triethylene tetramine (TETA) system applied on to aluminium by using both FTIR and XPS. They found close to the polymer/metal interface, the partial protonation of TETA by aluminium hydroxides. Kinzler et al. [57] used XPS to identify the chemical species present at the zinc oxide/organic layer interface.

CH3 (n + 1) HO

O

C

+

OH

(n+2) CH2

CH2

CH

Cl

+

(n + 2) NaOH

CH3

CH3

O CH2

CH

CH2

O

C

CH3 O

CH3

CH2

CH

CH2

OH

O

C

O O

CH3

+ (n + 2) NaCl + (n + 2) H2O

Scheme 15.4 Synthesis of DGEBA diepoxy prepolymers by the Taffy process [54].

CH2

CH

Cl

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Handbook of Multiphase Polymer Systems H R1

N

H2C

H

H

O

+

CH

R1

N

OH CH2

CH

(a) H R1

N

H2C

R2

R2

O

+

CH

R1

N

OH CH2

CH

(b) OH CH R2 R1

N

OH CH2

CH

R2

O

+ H2C

CH

R1

N

CH2 O

CH2

CH

(c)

Scheme 15.5 Chemical reactions for a diepoxy diamine system: (a) epoxide ring-opning reaction with primary amine; (b) epoxide ring-opening reaction with secondary amine; (c) etherfication reaction between reacted and unreacted epoxy groups.

XPS has been employed to investigate the adsorption of an adhesive and two of its major components – an amine curing agent (TDI urone) and DGEBA – on a hydrated aluminum surface and hydrated surface coating with GPS. The XPS data suggested donor-acceptor interactions between the curing agent and the substrate [58]. The interaction of diethanolamine (DEA), a model compound for an amine-cured epoxy resin, on airoxidized copper, aluminium and phosphoric acid anodized aluminium was studied using temperature programmed desorption (TPD) and XPS [59]. DEA desorption temperatures and N1s binding energies are consistent with N lone-pair coordination to acidic aluminium sites. Affrossman et al. [60] studied by XPS and static SIMS the interaction of DEA with aluminium surfaces and found N1s binding energy shifts consistent with a protonated state for DEA adsorbed on phosphoric acid anodized aluminium. For ion-bombarded aluminium surface, the N 1s signal showed a single binding energy in the range of Lewis site interactions on alumina. Elsewhere, several investigations were conducted on the relationship between substrate acidity and N1s binding energies of the adsorbates, notably by Borade et al. [61, 62] Other studies concerned the XPS study of adhesion of epoxy resins to polymer-coated silica [63] and metal oxides (MgO, Al2 O3 , SiO2 ) [64]. Attention was also paid to photodegradation [64], UV-accelerated degradation [65], thermal degradation [66] and hygrothermal ageing [67] of epoxy adhesives. Of relevance to the paint industry, high-resolution XPS and ToF-SIMS were used to investigate the segregation of minor components – a melamine–formaldehyde resin and a poly(acrylic) flow control agent – in a model epoxy resin [68]. XPS and ToF-SIMS have shown that the surfaces of these primers, which consist of complex mixture of organic components, are enriched with the melamine–formaldehyde resin. The conclusions are drawn from the C1s spectrum that could be peak fitted to account for all 11 carbon functionalities present in the three components of the organic coating. Combining XPS and ToF-SIMS permitted to show that when the flow control agent is excluded from the formulation, the surface of the paint film is enriched in the melamine–formaldehyde component. On addition of the flow control agent, such segregation is only identified by XPS; the ToF-SIMS spectrum resembles that of the flow control agent, an indication of monolayer segregation of this component at the paint/air interface. Such segregation phenomena

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XPS Studies of Multiphase Polymer Systems

611

are shown to be insensitive to substrate surface pretreatment. It is clear from this work that, although XPS and ToF-SIMS are both surface-sensitive techniques, the latter can detect compositional changes within the top monolayer. In the field of civil engineering, the use of epoxy adhesives is continuously growing. They are mainly applied to the repair of concrete structures by crack injection, or by adhesive bonding of composites. However, the mechanisms of adhesion between concrete and polymer adhesive are not well understood and thus the long-term stability of the concrete–adhesive joints is yet to be predicted. Epoxy resin DGEBA is one type of adhesive used in civil engineering for assembling concrete blocks and is crosslinked by primary amines such as DETA or triethylenetetramine (TETA) because this curing reaction occurs spontaneously at ambient temperature. In this context, there is a long-term collaboration between the authors of this chapter aimed at employing XPS to investigate the surface chemical composition of cement pastes before and after coating with fully formulated epoxy resin and their components taken separately. For the cement paste/hardener interface, inspection of the N1s peak at the cement/hardener interface indicated the existence of donor–acceptor interactions, where the aminated hardener plays the role of the Lewis base (donor) and the cement acts as the Lewis acid (Figure 15.15(a)) [69]. This also holds for a mixture of epoxy and DETA coated on a cement paste prepared from an ordinary Portland cement (OPC) at water/cement (W/C) ratio of 0.5. Interestingly, the tailed shape of the N1s implies peak fitting with an additional component at 402.2 eV (Figure 15.15(b)) that is assigned to interfacial protonation of the hardener [70]. The interactions of DETA at the cement paste/(epoxy+DETA) interface suggests strong adhesion due to the specific hydrogen bonding

×102

5500

N 1s

32

OPC-R5%+H1/3 st

Intensity (CPS)

30

5000 Intensity (CPS)

P1: JYS

28

26

free amine 4500

H bonding N+

4000

24

22 3500 412

408

404 400 Binding energy (eV) (a)

396

404

402 400 398 Binding energy (eV)

396

(b)

Figure 15.15 High-resolution N 1s spectra from (a) TETA hardener coated on a blended cement paste (1% wt.); and (b) a mixture of epoxy/DETA coated from CHCl3 at 5% wt. of epoxy on a hardened ordinary Portland cement paste; DETA is at 1/3rd stoichiometric ratio relative to DGEBA. (Figure 15.15(a) reprinted from [69]. Copyright (2002) with permission from John Wiley & Sons, Ltd., and 15.15(b) reprinted from [70]. Copyright (2008) with permission from John Wiley & Sons Ltd.)

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Handbook of Multiphase Polymer Systems C-C C-H

C-N

C-O

15k

12k carbonates

Intensity (a. u.)

P1: JYS

OPC-R5%+Hst

9k OPC-R5%+H2/3st

6k OPC-R5%+H1/3st

3k

K2p OPC

0 280

285

290

295

300

Binding energy (eV) Figure 15.16 High-resolution C1s spectra of OPC before and after coating by DGEBA/DETA mixtures where the resin is at 5 wt% relative to OPC and the hardener DETA (H) is in the stoichiometric or below the stoichiometric ratio.

and electrostatic interactions between a DETA and the mineral on the one hand, and crosslinking the epoxy resin on the other hand. The specific interfacial interaction of DETA from the (epoxy+DETA) mixture and the cement paste surface were confirmed by examination of the OH-stretching frequency shifts in the FTIR spectra of the mineral substrate. Moreover, similar IR shifts due to hydrogen bonding were observed for the model cement hydration product portlandite after coating by (epoxy+DETA) mixtures [70]. Inspection of the C1s fine structure from OPC coated by DGEBA/DETA systems (Figure 15.16) is interesting as it shows important changes after coatings. Two main C1s peaks are displayed by uncoated OPC, one at 285 eV due to surface organic contamination and one at 290 eV due to carbonates. Upon coating by DGEBA/DETA systems, there is a substantial increase of the relative C1s peak intensity around 286.5 eV corresponding to C-O bonds. This peak position indicates the detection of DGEBA. Moreover, upon addition of DETA in the formulations from sub-stoichiometric to stoichiometric molar ratio, the main peak centred at 285 eV broadens due to the addition of the hardener and therefore of C-N bonds the C1s spectral position of which is ∼ 285.7–286 eV. Moreover, there is strong supporting evidence for the crosslinking of epoxy at the cement surface. Indeed, the increase in the C1s component at 285.7 eV is due to crosslinking epoxy which results in oxirane ring opening. As shown in Scheme 15.5, it is clear that two C-O bonds (C1s at 286.5) are replaced by one C-O and one C-N (285.7 eV), hence the C1s peak broadening.

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XPS Studies of Multiphase Polymer Systems

O N H

N n

polyaniline (PANI)

613

O

n S

H polypyrrole (PPy)

n

polythiophene

Scheme 15.6 Chemical structure of PANI, PPy and PEDOT.

15.3.4

Conductive Polymers

Inherently conducting polymers (ICPs) have attracted a massive number of research groups worldwide for almost three decades [71–75]. The main ICPs are polyaniline (PANI), polyaromatics (e.g. polyparaphenylenes) and polyheterocycles such as polypyrrole (PPy), polythiophene and poly(ethylenedioxy)thiophene (PEDOT) of which chemical structures are shown in Scheme 15.6. For the purpose of this chapter, the emphasis will be on polypyrrole. This ICP can be synthesized electrochemically as thin films on a variety of electrodes such as platinum, ITO (indium doped titanium oxide), Au, glassy carbon [76], iron [77], 316L stainless steel of medical grade [78], copper [79] and carbon nanotubes [80] in the presence of a supporting electrolyte. It can also be chemically synthesized using FeCl3 or (NH4 )2 S2 O8 as oxidants, the chemical route having the advantage of yielding much larger scales of products. PPy chains bear positive charges that are counterbalanced by anions (e.g. Cl− , NO3 − , SO= 4 . . .), called ‘dopants’. The ratio of dopant-to-repeat unit is called the ‘doping level’. The nature and the extent of the counter anion are among the factors controlling the properties of the ICPs such as conductivity [81]. One can further take advantage of the chemical synthetic route to obtain colloidal particles [82–85], composites [86, 87], hollow spheres [85, 88], nanotubes [89] and nanocomposites [90, 91] for several purposes such as bioadsorbents [85, 92], HPLC stationary phases [93], chemical sensors [94], corrosion control by PPy-based coatings [95] and conductive fillers [96]. Because PPy is insoluble, stiff and inflexible and, moreover, is usually part of multicomponent materials, XPS is frequently used to examine the surface and interface chemical compositions of polypyrrole and to monitor surface chemical modification upon Br¨onsted acid-base treatment [97, 98], ozone-mediated oxidation [99], ion-exchange of dopant by anions from buffered aqueous solutions [100], degradation induced by ageing or thermal stress [101]. We and the groups of E. T. Kang (see for example the review published in 1993 [102]), in Singapore, and of Luigia Sabbatini in Italy [103] have spent time and effort in this regard. In the following, we will show the role of XPS in the study of the interfacial physicochemical properties of PPy materials. We will particularly address questions pertaining to the relationships between surface chemical composition and non-covalent (van der Waals and ionic interactions) and covalent interactions of polypyrrole. 15.3.4.1

London Forces and Dispersion of Polypyrrole in Polypropylene

It is well known that in the condensed phase, van der Waals interactions are dominated by London-induced dipole-induced dipole interactions. These interactions are ubiquitous and concern all materials without any exception. It is therefore important to determine the propensity of materials to develop London interactions, e.g. by contact angle measurements for flat surfaces [104] or by inverse gas chromatography for powders and fibers [105]. However, it is important to identify the nature and surface concentration of species or functional groups at the surface which induce changes in the London interaction energies. XPS measurements were related to the surface energy changes on several occasions. An illustration of such correlation concerns polypyrrole powders prepared in the presence of the surfactant sodium bis(2-ethylhexyl) sulfosuccinate

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Handbook of Multiphase Polymer Systems CH3 CH2

O

CH3CH2CH2CH2

CH

CH2

O

C

CH3CH2CH2CH2

CH

CH2

O

C

CH2

CH2

CHSO3 Na

O

CH3

Scheme 15.7 Chemical structure of AOT.

(known as aerosol OT or AOT, see structure in Scheme 15.7) at various initial pyrrole/surfactant molar ratios [106]. The submicrometer-sized powder particles were investigated in terms of surface chemical composition, surface energy, conductivity and dispersion in a polypropylene matrix. AOT was used in order to impart to the conductive polypyrrole a hydrophobic character that makes it a compatible conductive filler with the thermoplastic matrix polypropylene. Figure 15.17 shows the XPS survey scans and narrow C1s regions of a chemically-synthesized polypyrrole in the presence of the AOT surfactant and the reference AOT-free polypyrrole. The main features C1s, N1s, O1s, Cl2p and S2p are centred at 285, 400, 532, 198 and 168 eV, respectively. On addition of AOT (lower pyrrole/AOT ratio), the O1s/N1s and O1s/C1s intensity ratios increase, since AOT is containing oxygen in the ester and sulfonate groups (see chemical structure in Scheme 15.7). Similarly, C1s/N1s is increasing with AOT because N is an elemental marker for PPy. The increase in the ratio S2p/Cl2p intensity ratio is due to the addition of AOT and its insertion in PPy structure as a co-dopant with chlorides. The C1s spectra of PPyCl and PPyCl/AOT (7/1) have complex structures which originate from the convolution of PPy and AOT photoemissions. One can distinguish between the so-called α and β type carbon atoms in PPy, the C1s components of which are centred at 284.8 and 283.8 eV, respectively [107]. The complex structure at higher binding energy is due to the defects in the PPy backbone and surface oxidation. One can also note a low intensity feature around 291.5 eV assigned to a shake-up satellite that is due to the conjugated structure of the polymer backbone. When AOT is used in pyrrole polymerization, the AOT surfactant brings

C-C/C-H

C1s O1s N1s

PPy-Cl/AOT (7/1)

Cl2p S2p

I (a. u.)

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200

400 600 800 1000 Binding energy (eV)

COOR PPyCl/AOT (7/1) PPyCl

280

285 290 Binding energy (eV)

295

Figure 15.17 Survey and C1s high-resolution spectra of chloride-doped polypyrrole (PPyCl) and the same polymer prepared in the presence of AOT surfactant PPyCl/AOT (7/1) where the initial pyrrole/AOT molar ratio is 7/1.

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XPS Studies of Multiphase Polymer Systems

70

30 °C 50 °C

PPy-SO4

615

60

γSd (mJ/m²)

60 (7/1) (5/1) (4/1)

50 40

(2/1)

γSd (mJ/m²)

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40

30 3

6

9

12

15

Surface C/N ratio (a)

18

50

55

60

65

Contribution of C-C/C-H peak to C1s (%) (b)

Figure 15.18 IGC-determined γ Sd values versus (a) C/N ratio, and (b) contribution of C-C/C-H component to C1s peaks. (Figure 15.18(a) reproduced from Ref. [96].)

two major changes: (i) increase of the C-C/C-H fraction of the C1s region, and (ii) appearance of a peak at 289 eV due to the ester groups of the surfactant (see Table 15.2 for chemical shifts). The main peak at 285 eV has a strong contribution from the hydrophobic alkyl chains of the surfactant which results in the minimization of the surface energy of the powders as determined by inverse gas chromatography (IGC). Indeed, for an initial molar pyrrole/AOT of 2/1, the dispersive contribution to the surface energy (γ S d ) is as low as 36.6 mJ/m2 at 30◦ C, hence a polyethylene-like surface of the conductive polymer and its compatibility with polypropylene matrix. Figure 15.18 shows a relationship between γ S d and C/N ratios on the one hand (Figure 15.18(a)). On the other hand, γ S d is plotted versus the contribution of C-C/C-H component to the C1s peak area determined by peak fitting (Figure 15.18(b)). The approach shown in Figure 15.18(b) is more rigorous as it relates γ S d to a particular functional group that induces surface energy minimization and not merely all carbon atom types. Compared to the AOT-free PPy, this minimization of the surface energy by AOT was found to be beneficial for the dispersion of PPy particles in the PP matrix, PPy-PP adhesion and the formation of conductive paths in the thermoplastic PP matrix. In studies pertaining to thermoplastic polymer adsorption onto Cl-doped PPy, γ S d was found to decrease from 144 mJ/m2 at 48◦ C down to 33–37 mJ/m2 when PMMA was adsorbed from CCl4 ; the latter value being close to that of PMMA [108]. In considering all casting organic solvents, γ S d was related to the surface fraction of PMMA as estimated by XPS within the depth probed by the technique. Similar trends of surface energy minimization were observed for the competitive adsorption of PVC and PMMA onto PPy [109]. Elsewhere [110], for freshly synthesized polypyrrole powders of NO3 -doped PPy, XPS was employed to understand the changes in the surface energy of polypyrrole. Indeed, γ S d values were correlated with the doping level of the conductive polymer (NO3 /NPPy ratio; NPPy = nitrogen from PPy backbone) determined by XPS. For a doping level of about 17 and 33 %, γ S d values determined by IGC were found to be ∼52 and 115 mJ/m2 , respectively, at 48◦ C. 15.3.4.2

Acid-base Interactions and Thermoplastic Polymer Adsorption onto PPy

It is very well known that acid-base interactions are driving forces in adsorption, wetting and adhesion phenomena [111, 112]. In this regard, XPS was used to study the surface chemical composition changes upon

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Handbook of Multiphase Polymer Systems THF Dioxane Toluene CCl4

4

CH2 Cl2 DCM CHCl3

3

O1s/N1s

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2

1

0 -25

-20

-15

-10 DN

-5

0

5

10

15

20

25

AN

Figure 15.19 O1s/N1s intensity ratio versus donor numbers (DN) and acceptor number (AN) of basic and acidic solvents, respectively. Based on data reported in Ref. [113].

adsorption of PMMA onto chloride-doped polypyrrole (PPyCl) from neutral, acidic and basic solvents [113]. To account for the adsorption of PMMA onto PPyCl, O1s/N1s ratios were related to the Gutman’s donor and acceptor numbers (DN and AN, respectively) of the casting solvents. Nitrogen is a unique elemental marker for polypyrrole whilst oxygen has a major contribution from the PMMA adsorbate (O1s peak from neat PPy is due to slight surface oxidation). Figure 15.19 shows that the highest O1s/N1s ratios were obtained for adsorption of PMMA from non-competing solvents. In contrast, Lewis acidic solvents (e.g. chloroform) dissolve PMMA and prevent this polymer from adsorption on PPyCl, whereas Lewis bases (e.g. THF) compete for the acidic sites on PPyCl surface, thus hampering the adsorption of PMMA.

15.3.4.3

Ion Exchange Properties and Electrostatic Interactions of Polypyrrole

Polypyrrole has a positively charged backbone which is neutralized by the negatively-charged dopants. It is thus reasonable to expect favorable interactions with DNA fragments in buffered media. Saoudi et al. [98] have studied at various pH and ionic strengths the adsorption of DNA onto PPyCl by the depletion method using UV spectroscopy analysis of the supernatnant. DNA adsorption was found to be in the range of 0.13–0.55 mg/m2 with a maximum at low pH. Indeed, at high pH, polypyrrole is deprotonated and therefore loses its positively-charged character favorable for DNA adsorption. The deprotonation results in C=N imine species that can easily be detected by XPS on the N1s spectrum of polypyrrole [98]. The imine groups yield a component at 398.5 eV (see Table 15.2). XPS was further used ex,situ to determine DNA adsorption isotherms [100]. First, it was found that after overnight conditioning of polypyrrole in a phosphate buffer solution, chloride dopants were partially exchanged phosphates from the buffer. The survey Cl2p peak in the neat PPyCl was replaced by P2p

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O

Intensity (cps)

Intensity (cps)

617

970 C

1419

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O

N

582

388

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473

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PPyCl-NaP

0

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0 100 200 300 400 500 600 700 800 900 1000

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Binding energy (eV)

1365

Binding energy (eV)

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Na

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Intensity (cps)

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546

Intensity (cps)

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265 P

PPyCl-DNA 0.4 mg/m2 0

DNA 0

0 100 200 300 400 500 600 700 800 900 1000

Binding energy (eV)

0 100 200 300 400 500 600 700 800 900 1000

Binding energy (eV)

Figure 15.20 XPS survey scans of (a) PPyCl bulk powder, (b) PPyCl-NaP, (c) PPy-DNA (0.4 mg DNA/m2 ), (d) sodium salt DNA fibers. Reproduced from Ref. [100].

after buffer-conditioning (Figure 15.20). Ion-exchange chromatography permitted to detect a depletion of phosphates from the buffer and their concomitant replacement by chlorides leached from PPyCl [100]. After conditioning and incubation of PPy with DNA, the conductive polymer powders were analyzed by XPS and the surface composition reported in surface atomic ratios (Figure 15.21). XPS-determined X/N atomic ratios (X = O, P) were plotted versus DNA equilibrium concentration resulting in isotherms of high affinity type (Figure 15.21(a)).

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1.4

0.4

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O/N

1.0 0.8

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DNA equilibrium concentration (µg/ml)

0.24

0.36

0.48

0.60

DNA adsorption (mg/m2) PPyNO3

PPyCl

(a)

(b)

Figure 15.21 Adsorption of DNA on PPy powders at pH 7. (a) Isotherm plots of O/N and P/N versus equilibrium DNA concentration; (b) Na/P atomic ratios versus the amount of adsorbed DNA onto PPyCl and PPyNO3 . Reprinted from [96]. Copyright (2006) with permission from Elsevier.

At relatively high amounts of adsorbed DNA, both PPyCl and PPyNO3 surfaces exhibited increase in sodium cations adsorbed from the DNA solutions (Figure 15.21(b)). Increase in surface sodium content compensates for the excess of DNA negative charges at the extreme surface of the system. Indeed, the negative charges of DNA at the PPy-DNA interface are neutralized by positive charges of the polymer. 15.3.4.4

Covalent Bonding of Polypyrrole

Polypyrrole is known to be a strong adsorbent of molecules, polymers and biomacromolecules due to energetic interfacial van der Waals, electrostatic and hydrophobic interactions. However, it is also known that covalent linkages are preferred over non-covalent bonds if long-term stability of a joint is needed. Particularly, pyrrole can be copolymerized with N-succinimidyl ester-functionalized pyrrole (PyNSE) in order to obtain a reactive form of polypyrrole towards proteins (HSA) [84, 85] as schematically shown in Scheme 15.8. Such a synthesis was carried out by the in situ copolymerization of pyrrole 1 and the active PyNSE 2 with initial 1:2 fractions of 25:75 (%) in the presence of sterically stabilized polystyrene (PS) latex particles

PPyNSE

Py + Py-NSE FeCl3 PS

PS NSE

PS-PPyNSE

Scheme 15.8 Schematic synthesis of reactive PS-PPyNSE colloidal particles. Stars indicate surface reactive Nsuccinimidyl ester groups (NB: for clarity, size of polypyrrole coatings and NSE groups are out of scale).

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XPS Studies of Multiphase Polymer Systems

619

6

Γmax(HSA), mg/m2

5

4

3

2

1

0 0,31

0,32

0,33

0,34

0,35

Surface fraction of PyNSE, (O+N)/C

Figure 15.22 Maximal immobilization of HSA ( max ) versus the fraction of NSE groups at the surface of PSPPyNSEx particles, as determined by XPS. Reproduced from Ref. [114].

(600 ± 10 nm) [114, 115]. XPS indicated a substantial coating of PS by the reactive conducting copolymer due to the significant increase of N1s/C1s and O1s/C1s ratios on going from the sterically stabilized PS colloidal particles to the PS-PPyNSE colloids. Using a higher initial fraction of PyNSE gave a much higher surface (O+N)/C atomic ratios (determined by XPS) since PyNSE comonomer (C11 H12 N2 O4 ) bears much more oxygen and nitrogen atoms than pure pyrrole (C4 H5 N). The (O+N)/C atomic ratio was used as a chemical descriptor to account for the increase in PyNSE loading with increasing initial feed ratio. Moreover, this descriptor accounted for the maximal surface concentration of bound protein ( HSA(max) ). Figure 15.22 shows that  HSA(max) linearly increases with increasing surface fraction of Py-NSE repeat units which are responsible for covalent protein attachment.  HSA(max) increased by 25 %, when going from the minimal to the maximal surface functionnalization of the latex particles by NSE groups [114]. 15.3.5

Polymer Blends

Blending existing polymers to produce polymeric materials with new properties is a well-known method that is a simpler and economically more interesting alternative to polymerization of new monomers. Binary blends are subdivided into two categories: miscible and immiscible blends. As immiscible blends phase separate readily, it is expected that one of the component segregates to the surface and be detected by appropriate analytical tools such as XPS and ToF-SIMS. In bulk, phase separation can be probed by thermal analysis which shows two distinct glass transition temperatures (Tgs) corresponding to the individual polymers. Still, compatibilization of the two polymers can be achieved by mixing them with their corresponding block copolymers. Alternatively, it is possible to modify one of the polymers so that it can interact with the other through hydrogen bonding. This topic has been studied at length by Kwei and co-workers [116–118]. These authors have shown for several systems that

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one Tg is observed for blends interacting through hydrogen bonds. Moreover, they found Tgs with higher values than the weighed Tgs of the two individual polymers; the excess is being due to the interfacial hydrogen bonding [116, 119]. Although in bulk, several sets of polymers are known to be miscible (e.g. PMMA/PVC, PVP/PVPh) and to produce clear solutions and a single glass transition, using solid-state NMR, Kwei et al. [120] determined heterogeneities with a dimension between 25 and 2 nm for blends of syndiotactic PMMA and poly(styrene-covinylphenol). In a polymer mixture, a single composition-dependent glass transition is actually commonly taken as an indication of miscibility on the order of 20–40 nm [118]. Although polymers can be miscible in bulk, still there is possible surface segregation of one of the components which can be probed by XPS, as shown by several authors. Let’s start with a typical example of immiscible PS/PMMA symmetric blends. Kailas et al. [121] have shown by XPS surface enrichment by PS for solvent cast samples prepared from 50:50 (w/w) polymer mixtures (approximately 50:50 mol/mol). From the O and C at.% (5 and 95 %, respectively) and considering the stoichiometric formulae of PS and PMMA, the PMMA surface molar fraction was calculated using: % PMMA = (CPMMA /5)/[(CPMMA /5) + (CPS /8)] × 100% where 5 and 8 are the number of carbon atoms per MMA and styrene repeat units, respectively. The surface PMMA molar fraction was found to be 19.4% (and not 13% as reported in the paper because the authors did not take into consideration the number of carbon atoms per repeat MMA and styrene units). Upon annealing, the surface molar fraction of PMMA increased to 79.6% (and not 60% as reported in Ref. [121]). This XPS study was related to nano-SIMS mapping and AFM images, the latter techniques permitted to conclude that PS formed phase separated domains at the surface. Three dimensional nano-SIMS images suggested that the PS droplets were supported inside a rim of PMMA and these droplets continued from the surface like columnar rods into the film until the substrate–blend interface. Another interesting polymer blend with surface segregtion is that obtained from PS and poly(vinyl ethyl ether), PVEE. Beamson et al. [122] used ARXPS to investigate the surface and highlight segregation of PVEE to the surface. Figure 15.23 shows angle-resolved C1s spectra for a thin film of PS/PVEE blend sample. At 90◦ the C-C/C-H component (285 eV) dominates the spectrum indicating that below the surface, the blend is PS-rich (see Figure 15.3 for comparison), whereas at grazing angle, the C1s region quasi reduces to the resolved C1s doublet from pure PVEE (two carbon atoms in C-C/C-H bonds and two carbon atoms in C-O environments), implying that the latter segregates to the surface of the blend. The composition depth profiles CDPs were modeled by an empirical hyperbolic tangent function with φ b (PVEE bulk volume fraction),  (related to the surface overlayer thickness) and σ (width of interface between surface overlayer and bulk), as floating parameters. φ s (PVEE surface volume fraction) was assumed to be one. These are determined by nonlinear least squares regression, their uncertainties estimated and the curve fit residuals analyzed to demonstrate that the hyperbolic tangent CDP is a satisfactory fit to the ARXPS data. Conclusions are drawn regarding the behavior of the blend thin films as the thickness and polystyrene molecular weight are varied. Flory-Huggins interaction parameters (χ ) for the mixtures were calculated based upon the segregation data, and suggest a value of χ = 0.05 to be appropriate for the PS/PVEE system. Another well-known polymer blend is that of PMMA and PVC; it has been reviewed by Briggs [123]. PMMA and PVC can interact via Lewis acid-base interactions and miscibility is thus favored through such specific interactions. However, controversies were noticed concerning miscibility of the two polymers: the surface had the same composition as the bulk; segregation was noticed for PVC; PMMA segregated to the surface [123]. Such controversies could be due to the way the blends were prepared: choice of solvent or cosolvent, preparation by solvent casting or spin coating, annealing temperature. Using imaging XPS of the Cl2p and O1s regions (peak-background intensity), PMMA and PVC domains were probed (Figure 15.24) [124].

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35

× 103 θ = 90° θ = 75° θ = 60° θ = 45° θ = 30°

30

25

Intensity

621

θ = 20°

20

θ = 10° 15

10

5 0 292

290

288

286

284

282

280

Binding Energy (eV)

Figure 15.23 Angle-resolved C1s spectra for a PST (4.1 kDa)/PVEE (3.8 kDa) blend film 305 nm thick. Take-off angles are relative to the sample surface, therefore analysis at a grazing take-off angle of 10◦ reflects the outermost layers of the specimen. Reprinted from [100]. Copyright (2004) with permission from IOS Press.

The images are complementary (Figure 15.24) but the contrast is much greater in the Cl2p image. This was ascribed to a thin PMMA overlayer covering the domain structure extending almost to the surface. It is only with the help of imaging ToF-SIMS combined to sputtering that the imaging XPS results were confirmed. Indeed, sputtering the outermost layer permitted to collect more Cl– ions from PVC domains which were buried below a thin layer of PMMA covering the blend surface. These conclusions were also reached by Abel et al. [109] in their study of adsorption of PMMA/PVC mixtures from THF and dioxane on polypyrrole powder particles. XPS and ToF-SIMS showed that both polymers adsorbed on polypyrrole but that the underlying polypyrrole particles were less screened when the blend was coated from THF. Indeed ˚ when the blends were cast from THF and dioxane, the average overlayer thickness was less than 5 and 10 A respectively. Combining XPS, ToF-SIMS and IGC results suggested that PVC/PMMA blends adsorbed onto polypyrrole from THF may form homogeneous patches. However, PMMA is depleted to the free surface of the blend patches adsorbed from 1,4-dioxane (see Scheme 15.9). The polypyrrole-(PVC/PMMA) blend interface is thus PVC-rich at least in the case of 1,4-dioxane in agreement with ToF-SIMS results. Miscibility of polymers has also been studied at length in the case of hydrogen bonds between the mixed polymers. In this regard, a phenolic polysiloxane, poly(4-ethenylphenolmethylsiloxane) (PEPS), that contains the phenolic hydroxyl group as a hydrogen-bond donor has been synthesized and blended with a number of hydrogen-bond acceptors of different strengths, including poly(4-vinylpyridine) (PVPy) (strong), poly(vinylpyrrolidione) (PVPr) (strong), poly(dimethylacrylamide) (PDMA) (moderate), and poly(styreneco-acrylonitrile) (PSAN) (weak) [118]. All blends were miscible in the bulk, as indicated by a single DSC Tg, and were shown to be homogeneous by optical microscopy. XPS measurements demonstrated that all PEPS blends had surface enrichment in PEPS, which has a lower surface energy. Quantitative analysis of the

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Cl2p

O1s

Figure 15.24 XPS images of the surface of a PVC/PMMA film (cast from THF) based on the Cl2p and the O1s intensities (peak-background). The field of view is 700 μm. Reprinted from [100]. Copyright (2004) with permission from IOS Press.

sample surfaces was based on the relation: mole % of PEPS = (ISi2s /SSi2s )/[ISi2s /SSi2s + IN1s /SN1s ] where I Si2s and SSi2s are the integral peak intensity and sensitivity factor of Si2s, and I N1s and SN1s are the integral peak intensity and sensitivity factor of N1s. Figure 15.25 shows the relation between XPS-determined surface mole fraction of PEPS in the blends and its fraction in bulk. The linear plot is for equal bulk and surface compositions. Clearly, for all polymer blends the markers are well above the reference line, implying that PEPS surface segregates. This is driven by minimization of the surface energy (γ ) of the blends. This holds

PMMA-rich

PMMA-rich

PMMA-rich

PMMA / PVC

PMMA / PVC

PMMA / PVC

Polypyrrole

Scheme 15.9 Illustration of a polypyrrole surface partially coated with a PMMA/PVC blend.

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Surface mole fraction of PEPS

1.0

0.8

0.6 PVPy/PEPS PVPr/PEPS PDMA/PEPS PSAN/PEPS

0.4

0.2 0.2

0.3

0.4 0.5 0.6 Bulk mole fraction of PEPS

0.7

0.8

Figure 15.25 Surface versus bulk fraction of PEPS as determined by XPS for hydrogen bonding PEPS binary blends. Adapted from numerical data in Ref. [118].

for PVPy, PVPr and PDMA, that is higher segregation occurs for the largest difference between surface-free energies of the blended polymers (γ = 14.8, 22.4, 37.2, 60.8 and 32.2 mJ/m2 for PEPS, PVPy, PVPr, PDMA and PSAN, respectively). A reverse situation occurs with PSAN because the enthalpy of mixing for PEPS and PSAN was found to be the lowest and equal to –21kJ/mol, while it is –28.4, –24.6 and –23.9 kJ/mol for the systems PVPy/PEPS, PVPr/PEPS and PDMA/PEPS, respectively. The results indicated thus that (i) surface enrichment in hydrogen bonding polymer blends is governed by the interplay between differences in the surface energies of the constituents and the bulk thermodynamics; and (ii) hydrogen-bonding interactions reduce surface enrichment. Zeng et al. [119] studied hydrogen bonding in poly(vinyl phenol) (PVPh) and poly(4-vinyl pyridine) (PVPy) blends and complexes by time-of-flight secondary ion mass spectrometry (ToF-SIMS), XPS, and contact angle measurements. Systematic studies were performed on various blends of PVPh (Mw ∼30 kg/mol) and PVPy (Mw ∼ 150–200 kg/mol) in different solvents, including ethanol and N,N-dimethylformamide (DMF). Blends were prepared by mixing polymer solutions in DMF, whereas PVPh/PVPy complexes were precipitated from mixtures of polymer solutions in ethanol. All blends and complexes were found to be miscible with excess Tg due to hydrogen bonding as discussed above. Both XPS and wettability results showed no firm surface segregation of any component for the blends and complexes of PVPy and a low molecular weight PVPh. Excess of PVPh was found at the surface of the blends when a high molecular weight PVPh (Mw ∼300 kg/mol) was used. However, after annealing at 90◦ C, the surface and bulk compositions were found again to be the same, implying that the surface of blends of high molecular weight polymers may not be in the thermodynamic equilibrium state. The peak intensity of the characteristic pyridyl ions of the blends, especially the PVPh/PVPy complexes, such as the peak at m/z = 106 (pyridyl hydrogen bonded to OH groups from PVPh); was greatly enhanced due to hydrogen bonding. Figure 15.26 shows peak intensity ratio I106 /I107 versus the XPS-determined surface molar ratio PVPy/PVPh, where I106 is the peak corresponding to the positively-charged hydrogen bonded pyridyl and I107 that of the positively-charged + CH2 –C6 H4 –OH from PVPh. Only in the case of the blend with the mole ratio of PVPy:PVPh = 0:18; the intensities of

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5 4 I106/I107

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1 2 3 4 5 Surface mole ratio of PVPy/PVPh

6

Figure 15.26 I106 /I107 vs. XPS-determined surface mole ratio of PVPy/PVPh. Reprinted from [114]. Copyright (2004) with permission from Elsevier.

the peaks at m/z 106 and 107 are similar. It is suggested that the increase in the intensity of the peak at m/z = 106 is not caused by the surface segregation of PVPy, as confirmed by the results of XPS and contact angle measurements, but by the formation of hydrogen bonds between the pyridyl and hydroxyl groups. Polymer blends containing at least one (semi)crystalline polymer are interesting as the surface composition depends on the chemical nature of the blended polymers, their surface energies and cristallinity. In this regard, the polymers BA-C8 and 6FBA-C8 were obtained by condensation polymerization of 1,8-dibromooctane with bisphenol A and 4,4 -(hexafluoroisopropylidene)diphenol, respectively. XPS was used to analyze the surface of BA-C8-1/6FBA-C8 (80/20) blend as a function of crystallization time (BA-C8 is the semicrystalline polymer) [125]. Immediately after the sample was prepared, the surface concentration (computed from the F/O atomic ratio) of the fluorinated 6FBA was about 73 wt%, which is substantially higher than 20 wt% (bulk fraction), an indication of surface segregation despite bulk miscibility probed by one single glass transition temperature. After 800h, the surface concentration of 6FBA reached 94%. XPS results were consistent with those obtained by ToF-SIMS and contact angle measurements; the latter highlighting a shift to lower surface energy of the blend driven by 6FBA. 15.3.6

Composites

Perruchot et al. [126] reported the XPS study on polymer-grafted silica nanoparticles. This study indicated that the resulting inorganic–organic hybrid particles had core–shell morphologies. It demonstrates that initiatorfunctionalized silica particles can be used as a suitable colloidal substrate for the surface polymerization of non-ionic (OEGMA), cationic (MEMA) and anionic (SEM) hydrophilic methacrylates using aqueous ATRP. XPS permitted to highlight significant changes in the surface composition of silica nanoparticles by the initiator then by the polymer graft. In addition, the detection of unique elemental markers (i.e. N1s and S2p) (Figure 15.27) that are characteristic of the specific monomers, confirms the presence of thin polymeric grafted layers at the silica surface. Moreover, changes in the relative intensities of the peak-fitted carbon signals of the polymer-grafted silica particles compared to the reference silica particles provide further evidence for the grafted polymer chains. The C1s regions from the silica-polymer nanocomposites have structures typical of polymethacrylates, that is with a characteristic component at 289 eV due O-C=O functional group, a main peak at 285 eV and an intermediate one at 286.5 eV assigned to C-O bonds. The authors have also added for the peak fitting a component at ∼285.5 eV to account for the so-called beta shift mentioned above in Section 15.2.3. This study demonstrates that efficient surface-initiated ARTP of cationic, anionic and neutral hydrophilic monomers can be performed on colloidal silica particles and that the XPS characterized hydrophilic polymer grafts can efficiently disperse the inorganic silica particles.

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Figure 15.27 XPS survey spectra of (a) the bare silica sol; (b) the initiator-functionalized silica particles; (c) silicainit-POEGMA(500), (d) silica-init-PMEMA(500); and (e) silica-init-PSEM(500) particles, respectively. Reproduced from Ref. [126].

Djouani et al. [127] reported the XPS analysis of clay/PGMA nanocomposite synthesized by clay surfaceinitiated atom transfer radical polymerization (ATRP). The clay was first ion-exchanged with an ammonium compound bearing ATRP initiator functional group. Thermal gravimetric analysis (TGA) and XPS analyses indicated the nanocomposites have PGMA-rich bulk and surface, respectively. Indeed, the mass loading of PGMA reached 61 wt.% whilst XPS spectra, particularly the high resolution C1s region, resemble those of pure PGMA (Figure 15.28). Surface analysis permitted to understand the solubility of the nanocomposites

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I (cps)

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C-O

4k O-C=O 2k 0 280

285 Binding energy (eV)

290

Figure 15.28 (a) Survey spectra of MMT and MMT/PGMA nanocomposites; and (b) high resolution C1s region from MMT/PGMA. Reproduced from Ref. [127].

in chloroform. With such a high loading of PGMA, and the PGMA-rich surface of the MMT/PGMA nanocomposites, the latter were found to be fully compatible with epoxy resin without any sign of phase separation. The primary MMT/PGMA nanocomposites were mixed with DGEBA and DETA in chloroform in order to prepare molded, ternary MMT/PGMA–epoxy–DETA nanocomposites by solvent evaporation. The dried ternary systems exhibited superior viscoelastic properties (storage modulus and tan δ) compared to the neat crosslinked epoxy–DETA adhesive prepared in the absence of any clay nanocomposite. As far as conductive nanocomposites are concerned, montmorillonite/polypyrrole (MMT/PPy) nanocomposites, with 15% mass loading of PPy, were prepared by the in situ polymerization of pyrrole in the presence of montmorillonite (MMT) or organo-modified montmorillonite (oMMT) [91]. XPS showed that the MMT/PPy nanocomposite has an MMT-rich surface, whereas the oMMT/PPy nanocomposite surface has a rather organic nature. Due to the organic modification of MMT by the alkylammonium chloride, polymerization of pyrrole at the surface of oMMT is much more efficient in producing a conductive adlayer resulting in an enhancement of conductivity of the oMMT/PPy nanocomposites (1.1 S/cm) compared to MMT/PPy (3.1 × 10−2 S/cm). Figure 15.29 testifies indeed for relatively intense C1s and N1s peaks from the conductive polypyrrole coated on the organophillic clay (oMMT). The difference in the behavior of oMMT/PPy and MMT/PPy was interpreted in terms of surface energy minimization by the alkylammonium ions present at the surface of organo-modified MMT. Indeed, γ S d determined by IGC at 150◦ C, was estimated to be 34.0 mJ/m2 for oMMT, much lower than the value of 216 mJ/m2 determined for MMT. Firas Awaja [128] illustrated that XPS was able to explore the molecular and topographic changes of glass composites that are subjected to low earth orbit (LEO) environment and to provide the effects of atomic oxygen (AO), temperature cycling (TC), ultraviolet radiation (UV) and vacuum taken individually and in combination on the surface properties of a 3D glass. XPS results indicated the occurrence of chain scission and oxidation reactions on the surface of the treated samples. Samples that are subjected to AO generated more C-O and C=O functional groups on the surface. In addition, XPS indicated that all the treated samples had the largest oxygen concentration at the surface and the lowest surface carbon concentration. The sample that was treated with AO under vacuum showed the highest rate of chain scission and losses of volatiles in comparison with the other treated samples. Also, samples treated with AO showed the highest oxidation rate amongst the samples that were partially treated.

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1x106

Mg(KLL)

9x105

Intensity (a.u.)

Na1s

Fe2p

C1s

Si2p

8x105 7x105

627

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N1s

Al2p

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oMMT/PPy

4x105

oMMT

3x105 2x105 1x105

MMT

0 0

200

400

600

800

1000

Binding energy (eV) Figure 15.29

XPS survey regions of MMT, oMMT, MMT/PPy and oMMT/PPy. Reproduced from Ref. [91].

XPS had been used previously to study the changes to surface chemistry associated with treatment of epoxy resins. In a recent study, Awaja et al. [65] illustrated that XPS in conjunction with ToF-SIMS was able to reveal the surface degradation behavior of epoxy composites with different reinforcements and the molecular changes to carbon fiber, 3D and E-glass after intense UV and temperature conditions. It was shown that 3D glass exhibited a greater propensity to oxidize in comparison with the carbon fiber and E-glass samples. XPS measurements showed that all epoxy resin composites suffered significant surface oxidation. The 3D glass sample showed the most significant increase in oxygen concentration at the surface; approximately 50%. The E-glass, 3D glass and CF composites showed different relative ion intensities for aliphatic and oxygen-containing ions. Combining the ToF-SIMS positive and negative ion data with XPS data indicated the occurrence of the chemical phenomena of chain scission, crosslinking/condensation and oxidation as a result of the accelerated degradation. The change in the surface composition of the raw and chemically modified banana fiber was investigated by XPS [129]. Surface characterization by XPS showed the presence of numerous elements on the surface of the fiber. Investigation of the surface after alkali treatment, on the other hand, showed the removal of most of the elements. Silane treatment was found to introduce a considerable amount of silicon on the surface of the fiber. The O/C ratio was found to decrease in all cases except for the fluorinated and vinyl silane treated fibers. Detailed investigation of the deconvoluted C1s spectra revealed the change in the percentage atomic concentration of the various elements on the fiber surface. The XPS results were found to perfectly agree with the solvatochromic and electrokinetic measurements. 15.3.7

Interpenetrating Polymer Networks

An interpenetrated polymer network (IPN) is a combination of two polymers, in network form, of which at least one is synthesized and/or crosslinked in the immediate presence of the other without any covalent bonds

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Figure 15.30 XPS spectrum of a hybrid acrylate/zirconia IPN prepared on a silicon wafer showing the three characteristic peaks of carbon, oxygen and zirconium (PE = photoelectron emission). Note that the spectra were not calibrated for static charge as the C1s from the ester group of the acrylate is centred at 291 eV and not 289 eV. Reproduced from Ref. [131].

between them. These polymers are closely related to other multicomponent materials, containing completely entangled chains, such as polymer blends, grafts and blocks. However, the IPN can swell in solvents without dissolving and can suppress creep and flow. Most IPNs are heterogeneous systems comprised of one rubbery phase and one glassy phase which produce a synergistic effect yielding either high impact strength or reinforcement, both of which are dependent on phase continuity. There are four types of IPNs: sequential IPNs, simultaneous IPNs, semi-IPNs and homo-IPNs. XPS was used to investigate the complex polymerization of bismaleimide (BMI) in the presence of polyetherimide 1-methyl-2-pyrrolidinone (PEI/NMP) [130]. The surface chemical composition was found to depend on the use of ethanol and isopropyl alcohol as non-solvents. XPS permitted to highlight imide cleavage and oxa-Michael ethoxylation as concluded from FTIR analyses. Vendamme et al. [131] prepared and characterized hybrid nanofilm (35 nm-thick) of an organic acrylate network and an inorganic zirconia network. A typical XPS spectrum (Figure 15.30) of an IPN hybrid layer has three characteristic peaks of carbon, oxygen and zirconium. It was shown that the increase in the fraction of the inorganic precursor in the reactive formulation leads to an increased zirconium content in the nanofilm. Interestingly, the contents of carbon, oxygen and zirconium in the three different hybrids, as determined by XPS, essentially agreed with those calculated from the molar fraction of the different precursors in the initial formulation by assuming quantitative conversion of all of the monomers. In all cases, the agreement is excellent. It is clear that Zr(BuOn )4 precursor and the organic monomers were completely incorporated into the film. 15.3.8

Random and Block Copolymers

XPS has long been used to characterize the surface of copolymers [132]. For random copolymers XPS permits to highlight the presence of the comonomer repeat units in the end-copolymers [84, 88, 133, 134].

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This is possible if one of the monomers has a distinct element or a distinct functional group. Usually, surface analysis is related to initial feed ratio in order to understand the behavior of the polymeric material. Surface enrichment in one of the monomers constituting the random copolymers remains difficult to detect by XPS [134]. However, for some systems, such as membranes made from partially fluorinated poly(arylene ether sulfone) co- and terpolymers, surface analysis using XPS indicated that self-assembly occurred during membrane casting providing enhanced fluorine concentration on the surface of the membrane relative to the bulk calculations [135]. XPS study of block copolymers is interesting because these materials spontaneously exhibit various phase structures depending on the nature of the blocks, their respective length, the casting solvent, etc. In addition, the outermost surface can also be enriched in one of the blocks as proved by several authors and this since the pioneering paper of Clark et al. back in 1976 [136]. The surface enrichment in PDMS for the PDMS/PS block copolymers was found to depend on the casting solvent and ranged from 1.3 to 4 nm. Surface segregation of PDMS block was driven by surface energy difference between the blocks. When block copolymers form micelles, XPS permits to verify the core–shell structure and the shell nature induced by the aqueous or non-aqueous nature of the synthesis medium. In this regard, XPS was used to probe the nanostructure of shell crosslinked (SCL) micelles in the solid state, made from a MPC30 −GMA20 −DEA60 triblock copolymer (MPC: 2-methacryloyloxy phosphorylcholine; GMA: glycerol monomethacrylate; DEA: 2-(diethylamino)ethyl methacrylate) [137]. The latter comprised quaternized nitrogen atoms in the relatively hydrophilic MPC block, crosslinkable hydroxy groups in the middle GMA block, and neutral nitrogen atoms in the relatively hydrophobic DEA block. Two types of micelles were prepared: conventional micelles with the DEA block in the core in aqueous media and inverted micelles with the MPC block in the core in nonaqueous media. Efficient shell crosslinking of both types of micelles was achieved using divinyl sulfone (DVS), which reacts with the hydroxy groups in the central B block at room temperature. Although the XPS sampling depth of 2−10 nm is comparable to the dimensions of the micelle layers, analysis of the high-resolution N1s spectra provided the first direct evidence for the expected core−shell nature of SCL micelles in the solid state. In the absence of an elemental marker or characteristic functional group, the study of copolymers by XPS becomes too difficult. This is the case with polystyrene-polyethylene block copolymers. Turgeon and Paynter [138] showed that for such sp2-sp3 block copolymers the C KLL Auger line can bring useful information of the sp2/sp3 carbon atoms at the surface. Indeed, these authors defined a D-parameter that gives the separation in energy between the most positive and most negative excursions of the derivative (the D-parameter) plotted against the sp2 carbon concentration (in bulk). Figure 15.31 depicts an illustration of the determination of the D-parameter and how it correlates to the % of sp2 carbon atoms. Figure 15.31(a) shows the derivative of Auger C KLL regions from illustrates from polystyrene-polyethylene block copolymers and, graphite (inorganic sp2) and diamond (inorganic sp3) materials taken separately as well as their mixtures. In Figure 15.31(b), the D-parameter-%sp2 plot shows that the extent of sp2 carbon atoms could be deduced from the D-parameter. The authors clearly showed that, despite the presence of the shake-up satellite in the C1s region of PS at a BE well above that of the main peak (see Figure 15.3), this feature vanishes for sp2 carbon atom concentration less than 35%, hence the interest of the D-parameter.

15.4 Conclusion This chapter reviewed the basic principles of XPS and selected applications to polymeric systems. XPS is now routinely used in all aspects of polymer science and engineering. Its development over the last 40 years permitted to tackle the following aspects of polymers (and other materials): chemical structure, modification of polymer surfaces by physical and chemical methods, growth of thin coating, degradation, adhesion, adsorption, and covalent bonding, to name but a few studies.

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Figure 15.31 (a) Determination of D-parameter from first derivative C KLL spectra obtained from graphite, diamond and PS-PE block copolymers; and (b) its correlation with %sp2 carbon atoms in the same materials. Reproduced from ref. [138].

With the instrumental developments witnessed over the recent years (bright monochromatic sources, efficient detectors, high spectral resolution, imaging capability, etc.), XPS continues to attract researchers from diverse horizons for determining the surface chemistry of materials and thus understanding their macroscopic behavior. Combined with other tools (as clearly shown in this chapter), XPS plays a central role in research and development of materials with specific surface properties.

Glossary Abbreviation DEA DSPE-PEG DVS GMA MPC MWCNT PBS PεCL PDMA PDMS PE PEG PEI PEPS PES

Definition 2-(Diethylamino)ethyl methacrylate) 1,2-Distearoyl-sn-glycero-3-phosphoethanolamine-N-[methoxy (polyethylene glycol)-2000], a PEGylated phospholipids Divinyl sulfone Glycerol monomethacrylate 2-methacryloyloxy phosphorylcholine Multiwalled carbon nanotubes Poly(4-bromostyrene) P(ε-caprolactone) Poly(dimethylacrylamide) Poly(dimethyl siloxane) Polyethylene P(ethyelene glycol) Polyetherimide Poly(4-ethenylphenolmethylsiloxane) Poly(ether sulfone)

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PGMA PLGA PMEMA POEGMA PPy PPyNSE PS PSAN PSEM PSS PTFE PVC PVPh PVPr PVPy

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Poly(glycidyl methacrylate) Poly(lactic-co-glycolic acid) Poly(2-(N-morpholino)ethyl methacrylate) Poly(oligoethylene methacrylate) Polypyrrole N-succinimidyl ester-functionalized polypyrrole Polystyrene Poly(styrene-co-acrylonitrile) Poly(ammonium 2-sulfatoethyl methacrylate) Poly(styrene sulfonate) Poly (tetrafluoro ethylene) Poly(vinyl chloride) Poly(vinyl phenol) Poly(vinylpyrrolidione) Poly(4-vinyl pyridine)

References 1. NIST X-ray Photoelectron Spectroscopy Database. NIST Standard Reference Database 20, Version 3.5. Data compiled and evaluated by C. D. Wagner, A. V. Naumkin, A. Kraut-Vass, J. W. Allison, C. J. Powell, John R. Rumble, Jr. NIST Database. http://srdata.nist.gov/xps/ 2. A. A. Galuska and D. Eric Halverson, Quantitative analysis of surface ethylene concentrations in ethylene–propylene polymers using XPS valence bands, Surf. Interface Anal., 26, 425–432 (1998). 3. D. A. Shirley, C. S. Fadley, X-ray photoelectron spectroscopy in North America—the early years, J. Electron Spectrosc. Relat. Phenom., 137–140, 43–58 (2004). 4. G. Beamson and D. Briggs (Eds.) in High Resolution XPS of Organic Polymers. The Scienta ESCA300 Database. Wiley, Chichester (1992). 5. B. R. Strohmeier, The effects of O2 plasma treatments on the surface composition and wettability of cold-rolled aluminum foil, J. Vac. Sci. Technol. A, 7, 3238–3245 (1989). 6. M. P. Seah and W. A. Dench, Quantitative electron spectroscopy of surfaces: A standard data base for electron inelastic mean free paths in solids, Surf. Interface. Anal., 1, 2–11 (1979). 7. D. Briggs, M.P. Seah (Eds.), Practical Surface Analysis, Auger and X-Ray Photoelectron Spectroscopy, second ed., vol. 1, John Wiley, Chichester, UK, (1983) p. 201 p. 531. 8. D. Briggs (Ed.), Surface analysis of polymers by XPS and static SIMS. Cambridge University Press, Cambridge, U), 1998, pp. 34–39. 9. P. J. Cumpson, M. P. Seah, Elastic scattering corrections in AES and XPS. II. Estimating attenuation lengths and conditions required for their valid use in overlayer/substrate experiments, Surf. Interface Anal., 25, 430–446 (1997). 10. S. Tougaard, Surface nanostructure determination by x-ray photoemission spectroscopy peak shape analysis, J. Vac. Sci. Technol. A 14, 1415–1423 (1996). 11. S. Tougaard, Accuracy of the non-destructive surface nanostructure quantification technique based on analysis of the XPS or AES peak shape, Surf. Interface Anal., 26, 249–269 (1998). 12. http://www.quases.com/ 13. S. Tanuma, C. J. Powell, D. R. Penn. Surf. Interface. Anal. 21, 165–176 (1993). 14. S. Tougaard. Surf. Interface Anal. 25, 137–155 (1997). 15. S. Gam-Derouich, B. Carbonnier, M. Turmine, P. Lang, M. Jouini, D. Ben Hassen-Chehimi, M. M. Chehimi. Electrografted aryl diazonium initiators for surface-confined photopolymerization: a new approach to design functional polymer coatings. Langmuir, 26, 11830–11840 (2010).

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16. M. Lampim¨aki, K. Lahtonen, P. Jussila, M. Hirsim¨aki, M. Valden. Morphology and composition of nanoscale surface oxides on Fe–20Cr–18Ni{111} austenitic stainless steel. J. Electron Spectrosc. Relat. Phenom. 154, 69–78 (2007). / P. Jussila, H. Ali-L¨oytty, K. Lahtonen, M. Hirsim¨aki, M. Valden. Inhibition of initial surface oxidation by strongly bound hydroxyl species and Cr segregation: H2 O and O2 adsorption on Fe–17Cr. Surf. Sci. 603, 3005–3010 (2009) / I. Preda, A. Guti´errez, M. Abbate, F. Yubero, J. M´endez, L. Alvarez, L. Soriano. Interface effects in the Ni 2p x-ray photoelectron spectra of NiO thin films grown on oxide substrates. Phys. Rev. B 77, 075411-1 to 7 (2008). 17. V. Kisand, A. Kikas, E. Kukk, E. N˜ommiste, K. Kooser, T. K¨aa¨ mbre, R. Ruus, M. Valden, M. Hirsim¨aki, P. Jussila, M. Lampim¨aki, H. Aksela, S. Aksela. Substrate-induced effects in the creation and decay of potassium 2p core excitations in ultrathin films of KCl on copper. J. Phys.: Condens. Matter 20, 145206 (9pp) (2008). 18. M. Schleberger. Quantitative investigation of amorphous Fe/Ge and Fe/Si by inelastic peak shape analysis. Surf. Sci. 445, 71–79 (2000). 19. F. Yubero, C. Mansilla, F. J. Ferrer, J. P. Holgado, and A. R. Gonz´alez-Elipe. Size and shape of supported zirconia nanoparticles determined by x-ray photoelectron spectroscopy. J. Appl. Phys. 101, 124910 (2007). 20. C. L´opez-Santos, F. Yubero, J. Cotrino, A. Barranco, A. R. Gonz´alez-Elipe. Plasmas and atom beam activation of the surface of polymers. J. Phys. D: Appl. Phys. 41, 225209 (12pp) (2008) 21. S. Hajati, V. Zaporojtchenko, F. Faupel, S. Tougaard. Characterization of Au nano-cluster formation on and diffusion in polystyrene using XPS peak shape analysis. Surf. Sci. 601, 3261–3267 (2007). 22. M. C. L´opez-Santos, F. Yubero, J. P. Espin´os, A. R. Gonz´alez-Elipe. Non-destructive depth compositional profiles by XPS peak-shape analysis. Anal. Bioanal. Chem. 396, 2757–2768 (2010). 23. S. Hajati, S. Coultas, C. Blomfield, S. Tougaard, Nondestructive quantitative XPS imaging of depth distribution of atoms on the nanoscale, Surf. Interface Anal., 40, 688–691 (2008). 24. W. Geyer, V. Stadler, W. Eck, M. Zharnikov, A. G¨olzh¨auser, M. Grunze. Electron-induced crosslinking of aromatic self-assembled monolayers: Negative resists for nanolithography. Appl. Phys. Lett. 75, 2401–2403 (1999). 25. M. Steenackers, A. K¨uller, N. Ballav, M. Zharnikov, M. Grunze, R. Jordan. Morphology control of structured polymer brushes. Small 3, 1764–1773 (2007). 26. K. Artyushkova. Structure determination of nanocomposites through 3D imaging using laboratory XPS and multivariate analysis. J. Electron Spectrosc. Relat. Phenom. 178–179, 292–302 (2010). 27. L. J. Gerenser, XPS of in-situ plasma modified polymers. In Plasma surface modification of polymers, M. Strobel, C. S. Lyons, K. L. Mittal (Eds.) pp. 43–64. / Polymer surface modification, K. L. Mittal (Ed.), VSP, Utrecht (The Netherlands), 2000. / M. Manso , A. Rodriguez, A. Paul, J. Barrero, F. Rossi, Plasma functionalization, surface characterization and protein retention of multiple-sized polymer beads, Surf. Interface Anal., 38, 322–325 (2006). / D. J. Balazs, K. Triandafillu, Y. Chevolot, B.-O. Aronsson, H. Harms, P. Descouts, H. J. Mathieu, Surf. Interface Anal., 35, 301-309 (2003). / A. Vesel, M. Mozetic, A. Zalar, XPS characterization of PTFE after treatment with RF oxygen and nitrogen plasma, Surf. Interface Anal., 40, 661–663 (2008). 28. S. Edmondson, V. L. Osborne, W. T. S. Huck, Polymer brushes via surface-initiated polymerizations, Chem. Soc. Rev., 33, 14–22 (2004). 29. Y. Tsujii, K. Ohno, S. Yamamoto, A. Goto, T. Fukuda, Structure and properties of high-density polymer brushes prepared by surface-initiated living radical polymerization, Adv. Polym. Sci., 197, 1–45 (2006). 30. T. Matrab, M.M. Chehimi, C. Perruchot, A. Adenier, V. Guillez, M. Save, B. Charleux, E. Cabet-Deliry, J. Pinson, Novel approach for metallic surface-initiated atom transfer radical polymerization using electrografted initiators based on aryl diazonium salts, Langmuir, 21, 4686–4694 (2005). 31. Qin Jia Cai, Guo Dong Fu, Fu Rong Zhu, En-Tang Kang, Koon-Gee Neoh, GaAs-polymer hybrids formed by surface-initiated atom-transfer radical polymerization of methyl methacrylate, Ang. Chem. Int. Ed., 44, 1104–1107 (2005). 32. T. Matrab, M.M. Chehimi, J. Pinson, S. Slomkowski, T. Basinska, Growth of polymer brushes by atom transfer radical polymerization from glassy carbon modified by electro-grafted initiators based on aryl diazonium salts , Surf. Interface Anal., 38, 565–568, (2006). 33. M. N. Nguyen, T. Matrab, C. Badre, M. Turmine, M. M. Chehimi, Interfacial aspects of polymer brushes prepared on conductive substrates by aryl diazonium salt surface-initiated ATRP, Surf. Interface Anal., 40, 412–417 (2008).

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34. T. Matrab, M. Save, B. Charleux, J. Pinson, E. Cabet-Deliry, A. Adenier, M. M. Chehimi, M. Delamar, Grafting densely-packed poly(n-butyl methacrylate) chains from an iron substrate by aryl diazonium surface-initiated ATRP: XPS monitoring, Surf. Sci., 601, 2357–2366 (2007). 35. D. Li, Q. He, Y. Cui, J. Li, Thermoresponsive gold nanoparticles with adjustable lower critical solution temperature as colorimetric sensors for temperature, pH and salt Concentration, Chem. Mater., 19, 412–417 (2007). 36. F.J. Xu, S.P. Zhong, L.Y.L. Yung, Y.W. Tong, En-Tang Kang, K.G. Neoh, Thermoresponsive comb-shaped copolymer-Si(1 0 0) hybrids for accelerated temperature-dependent cell detachment, Biomater., 27, 1236–1245 (2006). 37. C. Sun, F. Zhou, L. Shi, B. Yu, P. Gao, J. Zhang, W. Liu, Tribological properties of chemically bonded polyimide films on silicon with polyglycidyl methacrylate brush as adhesive layer, Appl. Surf. Sci., 253, 1729–1735 (2006). 38. H. Ma, D. Li, X. Sheng, B. Zhao, A. Chilkoti, Protein-resistant polymer coatings on silicon oxide bysurface-initiated atom transfer radical polymerization, Langmuir, 22, 3751–3756 (2006). 39. B. Mrabet, M. N. Nguyen, A. Majbri, S. Mahouche, M. Turmine, A. Bakhrouf, M. M. Chehimi, Anti-fouling poly(2-hydoxyethyl methacrylate) surface coatings with specific bacteria recognition capabilities, Surf. Sci. 603, 2422–2429 (2009). 40. S. Gam-Derouich, M. N. Nguyen, A. Madani, N. Maouche, P. Lang, C. Perruchot, M. M. Chehimi. Aryl diazonium salt surface chemistry and ATRP for the preparation of molecularly imprinted polymer grafts on gold substrates. Surf. Interface Anal., 42, 1050–1056 (2010). 41. T. Matrab, M. N. Nguyen, S. Mahouche, P. Lang, C. Badre, M. Turmine, G. Girard, J. Bai, M. M. Chehimi, Aryl diazonium salts for carbon fiber surface-initiated atom transfer radical polymerization, J. Adhes., 84, 684–701 (2008). 42. T. Matrab, M. M. Chehimi, J. P. Boudou, F. Benedic, J. Wang, N.N. Naguib, J. A. Carlisle, Surface functionalization of ultrananocrystalline diamond using atom transfer radical polymerization (ATRP) initiated by electro-grafted aryldiazonium salts, Diamond Relat. Mater., 15, 639–644 (2006). 43. S. Mahouche Chergui, N. Abbas, T. Matrab, M. Turmine, E. Bon Nguyen, R. Losno, J. Pinson, M. M. Chehimi, Uptake of copper ions by carbon fiber/polymer hybrids prepared by tandem diazonium salt chemistry and in situ atom transfer radical polymerization, Carbon, 48, 2106–2111 (2010). 44. Y. Shaulov, R. Okner, Y. Levi, N. Tal, V. Gutkin, D. Mandler, A. J. Domb. Poly(methyl methacrylate) grafting onto stainless steel surfaces: application to drug-eluting stents. ACS Appl. Mater. Interfaces, 1, 2519–2528 (2009). 45. D. J. Dyer, Photoinitiated synthesis of grafted polymers, Adv. Polym. Sci., 197, 47–65, (2006). 46. (a) C. Slim, Y. Tran, M. M. Chehimi, N. Garraud, J.-P. Roger, C. Combellas, F. Kanoufi. Microelectrochemical patterning of surfaces with polymer brushes. Chem. Mater. 20, 6677–6685 (2008). / (b) F. Hauquier, T. Matrab, F. Kanoufi, C. Combellas. Local direct and indirect reduction of electrografted aryldiazonium/gold surfaces for polymer brushes patterning. Electrochim. Acta 54, 5127–5136 (2009). 47. M. C. Davies, R. A. P. Lynn, S. S. Davis, J. Hearn, J. F. Watts, J. C. Vickerman, A. J. Paul, Preparation of polymer latex particles with immobilized sugar residues and their surface characterization by X-ray photoelectron spectroscopy and time-of-flight secondary ion mass spectrometry, Langmuir, 9, 1637–1645 (1993) / M. C. Davies, R. A. P. Lynn, J. Hearn, A. J. Paul, J. C. Vickerman, and J. F. Watts, Surface chemical characterization using XPS and ToF-SIMS of latex particles prepared by the emulsion copolymerization of methacrylic acid and styrene, Langmuir, 12, 3866–3875 (1996). 48. Y. Deslandes, D. F. Mitchell, A. J. Paine, X-ray photoelectron spectroscopy and static time-of-flight secondary ion mass spectrometry study of dispersion polymerized polystyrene latexes, Langmuir, 9, 1468–1472 (1993). 49. E. Pacard, M. A. Brook, A. M. Ragheb, C. Pichot, C. Chaix, Elaboration of silica colloid/polymer hybrid support for oligonucleotide synthesis, Colloids Surf. B: Biointerfaces, 47, 176–188 (2006). 50. T. Basinska, S. Slomkowski, S. Kazmierski, A. Dworak, M. M. Chehimi, Studies of the surface layer structure and properties of poly(styrene/α-t-butoxy-ω-polyglycidol) microspheres by carbon nuclear magnetic resonance, X-ray photoelectron spectroscopy, and the adsorption of human serum albumin and γ -globulins, J. Polym. Sci. A: Polym. Chem., 42, 615–623 (2004). 51. R. D´ıaz-L´opez, N. Tsapis, D. Libong, P. Chaminade, C. Connan, M. M. Chehimi, R. Berti, N. Taulier , W. Urbach, V. Nicolas, E. Fattal, Phospholipid decoration of microcapsules containing perfluorooctyl bromide used as ultrasound contrast agents, Biomater., 30, 1462–1472 (2009).

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52. C. Pichot, Surface-functionalized latexes for biotechnological applications, Curr. Op. Colloid Interface Sci., 9, 213–221 (2004). 53. T. Basinska, S. Slomkowski, A. Dworak, I. Panchev, M. M. Chehimi, Synthesis and characterization of poly(styreneco-a-t-butoxy-o-vinylbenzyl-polyglycidol) microspheres, Colloid Polym. Sci., 279, 916–924 (2001). 54. P. Bardonnet (Ed.) R´esines Epoxydes : composants et propri´et´es. Techniques de l’ing´enieur (1992), trait´e Plastiques et Composites. 55. B. De’N`eve, Ph.D. Thesis, Ecole Nationale Sup´erieure des Mines de Paris, 1993. 56. R. G. Dillingham, F. J. Boerio, Interphase composition in aluminum/epoxy adhesive joints, J. Adhes., 24, 315–335 (1987). 57. M. Kinzler, D. M. Grunge, N. Blank, X-ray photoelectron microscopy applied to metal/epoxy laminates, J. Vac. Sci. Technol. A10, 2691–2697 (1992). 58. J. F. Watts, A. Rattana, M.-L. Abel, Interfacial chemistry of adhesives on hydrated aluminium and hydrated aluminium treated with an organosilane. Surf. Interface Anal., 36, 1449–1468 (2004). 59. J. A. Kelber and R. K. Brow. Model epoxy/metal-oxide chemical interactions: diethandamine on oxidized copper and aluminium. Appl. Surf. Sci., 59, 273–280 (1992). 60. S. Afrossman, R. F. Comrie, S. M. MacDonald. Interaction of a model epoxy resin compound, diethanolamine, with aluminium surfaces studied by static SIMS and XPS. J. Chem. Soc., Faraday Trans., 94, 289–294 (1998). 61. R. B. Borade, A. Sayari, A. Adnot and S. Kaliaguine, Characterization of acidity in ZSM-5 zeolites: an X-ray photoelectron and IR spectroscopy study, J. Phys. Chem., 94, 5989–5994 (1990). 62. R. B. Borade and A. Clearfeld, Characterization of acid sites in Beta and ZSM-20 zeolites, J. Phys. Chem., 96, 6729–6737 (1992). 63. J. H. Roh, J. H. Lee, N. I. Kim, H. M. Kang, T.-H. Yoon, K. H. Song, DSC analysis of epoxy molding compound with plasma polymer–coated silica fillers. J. Appl. Polym. Sci., 90, 2508–2516 (2003). 64. S. Scierka, A. Forster, W. Kosik, An XPS analysis of the interfacial interaction between oxides and epoxy resin. Federation of Societies for Coatings Technology, 492 Norristown Rd., Blue Bell, PA 19422 USA. 65. F. Awajaa, P. J. Pigram, Surface molecular characterisation of different epoxy resin composites subjected to UV accelerated degradation using XPS and ToF-SIMS, Polym. Degrad. Stab., 94, 651–658 (2009). 66. S.-G. Hong, The thermal-oxidative degradation of an epoxy adhesive on metal substrates: XPS and RAIR analyses, Polym. Degrad. Stab., 48, 211–218 (1995). 67. G. Z. Xiao, M. Delamar, M. E. R. Shanahan, Irreversible Interactions Between Water and DGEBA/DDA. Epoxy Resin During Hygrothermal Aging. J. Appl. Polym. Sci. 65, 449–458 (1997). 68. S. R. Leadley, J. F. Watts, C. J. Blomfield, C. Lowe. The Use of high-resolution XPS and ToF-SIMS to investigate segregation phenomena of minor components of a model coil coating formulation. Surf. Interface Anal. 26, 444–454 (1998). 69. M. I. Baeta Neves, V. Oliva, B. Mrabet, C. Connan, M. M. Chehimi, M. Delamar, S. Hutton, A. Roberts, K. Benzarti, Surface chemistry of cement pastes: a study by x-ray photoelectron spectroscopy, Surf. Interface Anal. 33, 834–841 (2002). 70. F. Djouani, C. Connan, M. M. Chehimi, K. Benzarti. Interfacial chemistry of epoxy adhesives on hydrated cement paste. Surf. Interface Anal. 40, 146–150 (2008). 71. T. A. Skotheim, (Ed.), Handbook of Conducting Polymers, vols 1 and 2, Marcel Dekker, New York (1986). 72. T. A. Skotheim, R. L. Elsenbaumer and J. R. Reynolds (Eds.), Handbook of conducting polymers, second edition. Marcel Dekker, New York (1998). 73. M. Aldissi (Ed.), Intrinsically conducting polymers: an emerging technology, Kluwer Academic Publishers, Dordrecht, The Netherlands (1993). 74. H. S. Nalwa (Ed.), Handbook of organic conductive molecules and polymers: Vol. 2. Conductive polymers: synthesis and electrical properties. John Wiley & Sons Inc, New York (1997). 75. T. A. Skotheim and J. Reynolds, Handbook of conducting polymers, third edition, CRC Press, Boca Raton, USA (2006). 76. S. Sadki, P. Schottland, N. Brodiec, G. Sabouraud, The mechanisms of pyrrole electropolymerization, Chem. Soc. Rev., 29, 283–293 (2000). 77. I. L. Lehra and S. B. Saidman, Morphology and properties of polypyrrole electrosynthesized onto iron from a surfactant solution, Synth. Met., 159, 1522–1528 (2009) / I. L. Lehra and S. B. Saidman, Corrosion protection

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78. 79. 80. 81.

82. 83.

84.

85.

86. 87. 88.

89. 90. 91.

92. 93. 94. 95.

96.

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of iron by polypyrrole coatings electrosynthesised from a surfactant solution, Corr. Sci., 49, 2210–2225 (2007). W. Khan, M. Kapoor, N. Kumar, Covalent attachment of proteins to functionalized polypyrrole-coated metallic surfaces for improved biocompatibility, Acta Biomater., 3, 541–549 (2007). M. Bazzaoui1, J. I. Martins, E. A. Bazzaoui, T. C. Reis, and L. Martins, Pyrrole electropolymerization on copper and brass in a single-step process from aqueous solution, J. Appl. Electrochem., 34, 815–822 (2004). S. Cosnier, M. Holzinger, Design of carbon nanotube-polymer frameworks by electropolymerization of SWCNTpyrrole derivatives, Electrochim. Acta, 53, 3948–3954 (2008). J. C. Thi´eblemont, M.F. Planche, C. Petrescu, J.M. Bouvier, G. Bidan, Stability of chemically synthesized polypyrrole films, Synth. Met., 59, 81–96 (1993) / K. Cheah, M. Forsyth, V.-T. Truong, C. Olsson-Jacques, Mechanisms governing the enhanced thermal stability of acid and base treated polypyrrole, Synth. Met., 84, 829–830 (1997). M. A. Khan, S. P. Armes, Conducting polymer-coated latexes, Adv. Mater., 12, 671 (2000). M. M. Chehimi, A. Azioune, S. Bousalem, A. Ben Slimane, A. Yassar, Preparation, characterisation and biomedical applications of conducting polymer particles, in Colloid polymers. Synthesis and characterization, H. Elaissari (Ed.), Marcel Dekker Inc. (NY), Surfactant Science Series, 115, 245–284 (2003), Chap.10. A. Azioune, A. Ben Slimane, L. Ait Hamou, A. Pleuvy, M. M. Chehimi, C. Perruchot, S. P. Armes, Synthesis and characterization of active ester-functionalized polypyrrole-silica nanoparticles: application to the covalent attachment of proteins, Langmuir, 20, 3350–3356 (2004). S. Bousalem, S. Benabderrahmane, Y. Yip Cheung Sang, C. Mangeney, M. M. Chehimi, Covalent immobilization of human serum albumin onto reactive polypyrrole-coated polystyrene latex particles, J. Mater. Chem., 15, 3109–3116 (2005). M. Omastov´a, K. Boukerma, M. M. Chehimi, M. Trchova, Novel silicon carbide/polypyrrole composites; preparation and physicochemical properties, Mater. Res. Bull., 40, 749–765 (2005). A. Ben Slimane, C. Connan, M.-J. Vaulay, M. M. Chehimi, Preparation and surface analysis of pigment@polypyrrole composites, Colloids Surf. A: Physicochem. Eng. Asp., 332, 157–163 (2009). C. Mangeney, S. Bousalem, C. Connan, M.-J. Vaulay, S. Bernard, M. M. Chehimi, Latex and hollow particles of reactive polypyrrole: preparation, properties, and decoration by gold nanospheres Langmuir, 22, 10163–10169 (2006). / A. Madani, B. Nessark, R. Brayner, H. Elaissari, M. Jouini, C. Mangeney, M. M. Chehimi. Carboxylic acid-functionalized, coreeshell polystyrene@polypyrrole microspheres as platforms for the attachment of CdS nanoparticles. Polymer, 51, 2825–2835 (2010). J. Jang, H. Yoon, Formation mechanism of conducting polypyrrole nanotubes in reverse micelle systems, Langmuir, 21, 11484–11489 (2005). K. Boukerma, J.-Y. Piquemal, M. M. Chehimi, M. Mravˇca´ kov´a, M. Omastov´a, P. Beaunier, Synthesis and interfacial properties of montmorillonite/polypyrrole nanocomposites, Polymer, 47, 569–576 (2006). M. Mravˇca´ kov´a, K. Boukerma, M. Omastov´a and M. M. Chehimi, Montmorillonite/polypyrrole nanocomposites. The effect of organic modification of clay on the chemical and electrical properties, Mater. Sci. Eng. C, 26, 306–313 (2006). S. Cosnier, Affinity biosensors based on electropolymerized films, Electroanal. 17, 1701–1715 (2005). C. Perruchot, M. M. Chehimi, M. Delamar, F. Dardoize, C haracterisation of the chromatographic properties of a silica–polypyrrole composite stationary phase by inverse liquid chromatography, J. Chrom. A, 969, 167–180 (2002). U. Lange, N. V. Roznyatovskaya1, V. M. Mirsky, Conducting polymers in chemical sensors and arrays, Anal. Chim. Acta, 614, 1–26 (2008). D.M. Lenz, M. Delamar, C.A. Ferreira, Application of polypyrrole/TiO2 composite films as corrosion protection of mild steel, J. Electroanal. Chem., 540, 35–44 (2003). / D. M. Lenz, M. Delamar, C. A. Ferreira, Improvement of the anticorrosion properties of polypyrrole by zinc phosphate pigment incorporation, Prog. Org. Coat., 58, 64–69 (2007). / D.E. Tallman, G.P. Bierwagen, Corrosion protection using conducting polymers, in: Handbook of conducting polymers, third edition: Conjugated polymers - processing and applications, CRC Press, 2006, pp. 15–1 to 15–53 (Chapter 15). K. Boukerma, M. Miˇcuˇs´ık, M. Mravˇca´ kov´a, M. Omastov´a, M.-J. Vaulay, P. Beaunier, M. M. Chehimi, Surfactantassisted control of the surface energy and interfacial molecular interactions of polypyrrole, Colloids Surf. A Physicochem. Eng. Asp., 293, 28–38 (2007).

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97. E. T. Kang, K. G. Neoh and K. L. Tan, The intrinsic redox states in polypyrrole and polyaniline: A comparative study by XPS, Surf. Interface Anal., 19, 33–37 (1992). 98. B. Saoudi, N. Jammul, M.-L. Abel, M. M. Chehimi and G. Dodin, DNA adsorption onto conducting polypyrrole, Synth. Met., 87, 97–103 (1997). 99. E. T. Kang, K. G. Neoh, X. Zhang, K. L. Tan and D. J. Liaw, Surface modification of electroactive polymer films by ozone treatment, Surf. Interface Anal., 24, 51–58 (1996). 100. B. Saoudi, N. Jammul, M. M. Chehimi, A.-S. Jaubert, C. Arkam and Michel Delamar, XPS study of the adsorption mechanisms of DNA onto polypyrrole particles, Spectrosc. Int. J., 18, 519–535 (2004). 101. M. M. Chehimi, E. Abdeljalil, A study of the degradation and stability of polypyrrole by inverse gas chromatography, X-ray photoelectron spectroscopy, and conductivity measurements, Synth. Met., 145, 15–22 (2004). 102. E. T. Kang, K. G. Neoh and K. L. Tan, X-ray photoelectron spectroscopic studies of electroactive polymers, Adv. Polym. Sci., 106, 135–190 (1993). 103. L. Sabbatini, C. Malitesta, E. De Giglio, I. Losito, L. Torsi, P. G. Zambonin, Electrosynthesised thin polymer films: the role of XPS in the design of application oriented innovative materials, J. Electron Spectrosc.Relat. Phenom., 100, 35–53 (1999). 104. A. Azioune, M. M. Chehimi, B. Miksa, T. Basinska, S. Slomkowski, Hydrophobic protein-polypyrrole interactions: the role of van der Waals and Lewis acid-base forces as determined by contact angle measurements, Langmuir, 18, 1150–1156 (2002). 105. A. Voelkel, B. Strzemiecka, K. Adamska, K. Milczewska, Inverse gas chromatography as a source of physiochemical data, J. Chrom. A, 1216, 1551–1566 (2009). 106. K. Boukerma, M. Omastov´a, P. Fedorko, M. M. Chehimi, Surface properties and conductivity of bis(2-ethylhexyl) sulfosuccinate-containing polypyrrole, Appl. Surf. Sci. 249, 303–314 (2005). 107. C. Malitesta, L. Losito, L. Sabbatini, P.G. Zambonin, New findings on polypyrrole chemical structure by XPS coupled to chemical derivatization labelling, J. Electron Spectrosc. Relat. Phenom., 76, 629–634 (1995). 108. M. M. Chehimi, M-L. Abel and Z. Sahraoui, An inverse gas chromatographic study of the PMMA/conducting polypyrrole interface, J. Adhes. Sci. Technol., 10, 287–303 (1996). 109. M.-L. Abel, M. M. Chehimi, F. Fricker and M. Delamar, A. M. Brown and J. F. Watts, Adsorption of PMMA and PVC blends onto polypyrrole: a study by XPS, ToF-SSIMS and inverse gas chromatography, J. Chrom. A, 969, 273–285 (2002). 110. M. M. Chehimi, M.-L. Abel, C. Perruchot, M. Delamar, S. F. Lascelles, S. P. Armes, The determination of the surface energy of conducting polymers by inverse gas chromatography at infinite dilution, Synth. Met., 104, 51–59 (1999). 111. F. M. Fowkes, Quantitative characterization of the acid-base properties of solvents, polymers, and inorganic surfaces, J. Adhes. Sci. Technol., 4, 669–691 (1990) 112. M. M. Chehimi, A. Azioune, E. Cabet-Deliry, Acid-base interactions: relevance to adhesion and adhesive bonding, in Handbook of adhesive technology, K. L. Mittal and A. Pizzi (Eds.), Marcel Dekker Inc. (NY), 2003, Chap. 5, pp. 95–144. 113. M.-L. Abel, M. M Chehimi, Effect of acid-base interactions on the adsorption of PMMA on chloride-doped polypyrrole from neutral, acidic and basic solvents: an XPS study. Synth. Met., 66, 225–233 (1994). 114. S. Bousalem, C. Mangeney, Y. Alcote, M. M. Chehimi, T. Basinska, S. Slomkowski, Immobilization of proteins onto novel, reactive polypyrrole-coated polystyrene latex particles, Colloids Surf. A Physicochem. Eng. Asp., 249, 91–94 (2004). 115. S. Benabderrahmane, S. Bousalem, C. Mangeney, A. Azioune, M.-J. Vaulay, M. M. Chehimi, Interfacial physicochemical properties of functionalized conducting polypyrrole particles, Polymer, 46, 1339–1346 (2005). 116. T. K. Kwei, The effect of hydrogen bonding on the glass transition temperatures of polymer mixtures, J. Polym. Sci. Polym. Lett. Ed. 22, 307–313 (1984) / T. K. Kwei, E. M. Pearce, J. R. Pennacchia, M. Charton, Correlation between the glass transition temperatures of polymer mixtures and intermolecular force parameters, Macromolecules, 20, 1174–1176 (1987). 117. E. M. Pearce, T. K. Kwei, S. Lu, Miscible blends through hydrogen bonding: effects on polymer properties, Polym. Adv. Technol. 5, 600–602 (1994).

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118. Y. Duan, E. M. Pearce, T. K. Kwei, X. Hu, M. Rafailovich, J. Sokolov, K. Zhou, S. Schwarz, Surface enrichment in polymer blends involving hydrogen bonding, Macromolecules 34, 6761–6767 (2001) 119. X.-M. Zeng, C.-M. Chan, L.-T. Weng, L. Li, Surface characterization and quantitative study of poly(4-vinyl phenol) and poly(4-vinyl pyridine) blends by XPS and ToF-SIMS, Polymer 41, 8321–8329 (2000). 120. L. Jong, E. M. Pearce, T. K. Kwei, NMR study of hydrogen bonded polymer blends: influence of the tacticity of poly(methyl methacrylate) on its miscibility with poly(styrene-co-vinylphenol), Polymer, 34, 48–55 (1993). 121. L. Kailas, B. Nysten, J.-N. Audinot, H.-N. Migeon, P. Bertrand, Multitechnique characterization of thin films of immiscible polymer systems: PS-b-PMMA diblock copolymers and PS-PMMA symmetric blends, Surf. Interface Anal., 37, 435–443 (2005) 122. G. Beamson, P. Mokarian-Tabari, M. Geoghegan, Composition depth profiling of polystyrene/poly(vinyl ethyl ether) blend thin films by angle resolved XPS, J. Electron Spectrosc. Relat. Phenom. 171, 57–63 (2009). 123. D. Briggs in ref. 8, section 6.5, pp. 177–184. 124. D. Briggs, Spectrosc.Eur. 5, 8 (1993). 125. Y.G. Lei, Z.-L. Cheung, K.-M. Ng, Lin Li, Lu-Tao Weng, Chi-Ming Chan, Surface chemical and morphological properties of a blend containing semi-crystalline and amorphous polymers studied with ToF-SIMS, XPS and AFM, Polymer 44, 2883–2890 (2003). 126. C. Perruchot, M. A. Khan, A. Kamitsi, S. P. Armes, J. F. Watts, T. von Werne, T. E. Patten. XPS characterisation of core–shell silica–polymer composite particles synthesised by atom transfer radical polymerisation in aqueous media. Eur. Polym. J. 40, 2129–2141 (2004). 127. F. Djouani, F. Herbst, M.M. Chehimi, K. Benzarti. Synthesis, characterization and reinforcing properties of novel, reactive clay/poly(glycidyl methacrylate) nanocomposites. Construct. Build. Mater. 25, 424–431 (2011). 128. Firas Awaja, Jin Bum Moon, Shengnan Zhang, M. Gilbert, Chun Gon Kim, P. J. Pigram. Surface molecular degradation of 3D glass polymer composite under low earth orbit simulated space environment. Polym. Degrad. Stab., 95, 987–996 (2010). 129. L. A. Pothan, F. Simon, S. Spange, S. Thomas. XPS Studies of chemically modified banana fibers. Biomacromolecules, 7, 892–898 (2006). 130. J. Kurdi, A. Kumar, Structuring and characterization of a novel highly microporous PEI/BMI semi-interpenetrating polymer network. Polymer, 46, 6910–6922 (2006). 131. R. Vendamme, S.-Y. Onoue, A. Nakao, T. Kunitake, Robust free-standing nanomembranes of organic/inorganic interpenetrating networks. Nature Mater. 5, 494–501 (2006) 132. F. Garbassi, M. Morra, E. Occhiello (Eds.), Polymer surfaces. From physics to technology. John Wiley & Sons Ltd, Chichester 1994, p. 103. 133. Shiyong Liu, Lu-Tao Weng, Chi-Ming Chan, Lin Li, Nick K. Ho, Ming Jiang. Quantitative surface characterization of poly(styrene)/poly(4-vinyl phenol) random and block copolymers by ToF-SIMS and XPS. Surf. Interface Anal. 31, 745–75 (2001) / A. Sezai Sarac, S. A. M. Tofail, M. Serantoni, J. Henry, V. J. Cunnane, J. B. McMonagle, Surface characterisation of electrografted random poly[carbazole-co-3-methylthiophene] copolymers on carbon fiber: XPS, AFM and Raman spectroscopy. Appl. Surf. Sci. 222, 148–165 (2004). 134. Tatsuo Teraya, Atsushi Takahara, Tisato Kajiyama, Surface chemical composition and surface molecular mobility of diblock and random copolymers with hydrophobic and hydrophilic segments. Polymer 31, 1149–1153 (1990). 135. K.B. Wiles, C.M. de Diego, J. de Abajo, J.E. McGrath. Directly copolymerized partially fluorinated disulfonated poly(arylene ether sulfone) random copolymers for PEM fuel cell systems: Synthesis, fabrication and characterization of membranes and membrane–electrode assemblies for fuel cell applications. J. Membrane Sci. 294, 22–29 (2007). 136. D. T. Clark, J. Peeling, J. M. O’Malley. Application of ESCA to polymer chemistry .8. surface-structures of ab block copolymers of polydimethylsiloxane and polystyrene. J. Polym. Sci. Polym. Chem. Ed. 14, 543–551 (1976). 137. Shiyong Liu, Yinghua Ma, S. P. Armes, C. Perruchot, J. F. Watts. Direct verification of the core–shell structure of shell cross-linked micelles in the solid state using x-ray photoelectron spectroscopy. Langmuir, 18, 7780–7784 (2002). 138. S. Turgeon, R. W. Paynter, On the determination of carbon sp2 / sp3 ratios polystyrene-polyethylene copolymers by photoelectron spectroscopy. Thin Solid Films, 394, 44–48 (2001).

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16 Light Scattering Studies of Multiphase Polymer Systems Yajiang Huang, Xia Liao, Qi Yang, and Guangxian Li College of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Sichuan, China

16.1 Introduction When a matter exposes under light, the electric field of the light induces an oscillating polarization of electrons in the molecules. Therefore, the molecules provide a secondary source of light and subsequently scatter light. The frequency shifts, the angular distribution, the polarization, and the intensity of the scattering light are determined by the size, shape and molecular interactions in the scattering material. Based on the pioneer work of Peter Debye [1], light scattering technology is utilized to characterize the microstructure of polymers [2], such as molecular weight and mean square radius of gyration. In the 1960s, Stein et al. [3] developed the light scattering theory for solids, and applied the modified theory to investigate the structure of polymer crystal. Static and dynamic light scattering are two experimental methods for measuring the patterns of scattered light. Static light scattering (SLS) (also known as classic light scattering or elastic light scattering) measures light intensity as a function of scattering angle and sample concentration. This allows the determination of average molecular weight, radius of gyration, and shape information. Dynamic light scattering (DLS) (also known as photon correlation spectroscopy or quasi-elastic light scattering) measures the time-dependent fluctuations in the intensity of scattered light [4, 5]. Analysis of the intensity fluctuations allows the determination of the diffusion coefficients, distribution of the particles, which can be converted into a size distribution using established theories. The molecular structure, conformation and orientation of polymer molecules greatly affect macroscopic properties of the material. The phase morphology and phase transition in multiphase systems play an important role in the potential application of the resulting material. Conventional characterization methods such as microscopy and differential scanning calorimetry (DSC) are widely used to investigate the morphology and miscibility of polymer blends under thermodynamic equilibrium conditions. However, these techniques

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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are normally destructive or discrete, and thus often used for offline measurement. Light scattering is a nondestructive technique with rapid response to density and orientation fluctuation of multiphase polymer systems; therefore, it has been widely utilized not only in the study of the phase behavior, crystallization and gelation of multiphase polymer systems under static conditions, but also in the in situ monitoring of the change of morphology of multiphase polymer systems under shear flow or in processing conditions. In this chapter, the experimental setup and theoretical framework, and the method of data analysis of light scattering has been described. The phase diagram of polymer blends and the possible mechanisms for phase separation, nucleation and growth, spinodal decomposition, and percolation-cluster transition, studied by light scattering technique have been discussed. Moreover, the application of time resolved-SALS (TR-SALS) in studying the reaction-induced phase-separation behavior and the structure development of thermoset/thermoplastics blends were summarized. In addition to the samples in equilibrium conditions, the application of light scattering for the in situ investigation of the phase separation dynamic and morphology development of polymer blends under shear flow, during extrusion or mixing were also discussed. The research into light scattering on the gelation process and crystallization of polymer blends has also been presented. The inhomogeneity of gels has been observed in both chemical gels and thermoreversible physical gels. There are four types of inhomogeneities in the gel state and four methods for the determination of the gelation threshold by considering the loss of ergodicity during the gelation process. The dependence of the crystalline morphology development and crystallization behavior in binary crystalline-amorphous or crystalline-crystalline polymer blends on the interlay of phase separation and crystalline has been discussed.

16.2 Light Scattering Technique The scattering phenomena in heterogeneous systems originate from the fluctuation in the spatial density (refraction index) or orientation of systems. The resolution of light scattering depends on the light wavelength (λ) of light source utilized. Usually, red laser light with λ = 632.8 nm or green laser light with λ = 532 nm are used, which correspond to a detective length scale ranging from 10−1 μm to 102 μm. Fortunately, many structures formed via phase separation, crystallization and gelation et al., which determine the optical and mechanical properties of final products in multiphase polymer systems, fall in this range. A good correlation between the detective resolution of light scattering and the dimension of the microstructure explored enables the successful employment of the light scattering technique in revealing the structure–property relationship of materials. 16.2.1

Scattering from Multiphase Polymer Systems

According to the electromagnetic field theory, as an incident visible light I0 propagates through a martial, the scattering by material (Rayleigh scattering) can be represented as [2]: Is =

c 2 α 2 ω4 E sin2 ϕ 4π 0 c4r 2

(16.1)

where ω is the frequency of light, E 0 the amplitude of electromagnetic field, α the polarity. The scattering function for a scattering body with spherical symmetry is: Is = 4π K  Vs η¯ 2



∞ 0

γ (r )

sin(qr ) 2 r dr qr

(16.2)

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where q is the scattering vector defined as q = 4π sin θ2 , Vs , λ and θ are the volume of scattering body, λ wavelength and the scattering angle measured in the medium, respectively. The scattering intensity is proportional to the mean square fluctuation of system η¯ 2 and also depends on the correlation function γ (r ). For a homogenous system there is no scattering and η¯ 2 = 0. By determining the angular dependence of scattering, the correlation function can be obtained via inverse Fourier transformation of I (q): γ (r ) =

c η¯ 2



∞

I (q) 0

sin(qr ) 2 q dq qr

(16.3)

According to Eq. (16.3), when r = 0, γ (0) = 1, thus: η¯ 2 = c



∞

I (q)q 2 dq

(16.4)

0

The integration in Eq. (16.4) is usually defined as invariant Q,  Q=

∞

I (q)q 2 dq

(16.5)

0

By determining I (q) experimentally, the invariant Q and mean square fluctuation η¯ 2 can be obtained without knowing the exact form of correlation function. Usually the correlation function has an exponential form: γ (r ) = e−r/αc

(16.6)

where αc is the correlation distance, which is the average distance of the fluctuation. In these cases, the detailed structure of the system can be characterized by the η¯ 2 and the correlation distance αc . The η¯ 2 and αc also influence the scattering intensity of the material at different scattering angles. For a two-phase mixture consisting of one domain with polarity of a1 dispersed in another domain with polarity of a2 , the average polarity of the mixture can be represented as: α¯ = φ1 α1 + φ2 α2

(16.7)

where φ1 and φ2 are the volume fraction of phases with φ1 + φ2 = 1. Thus the fluctuation of polarity in the two phases is: η1 = α1 − α¯ = −φ2 (α2 − α1 ) , η2 = α2 − α¯ = φ1 (α2 − α1 )

(16.8)

And the average fluctuation of system is: η¯ 2 = φ1 η12 + φ2 η22 = φ1 φ2 (α2 − α1 )2

(16.9)

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Namely, the average fluctuation is proportional to the multiple of volume fraction of the two phases and the square of difference of polarity between the two phases. Combining Eq. (16.2) with Eq. (16.6), we have: Is (q) = K



 0

∞

e−r/αc

sin(qr ) 2 αc3 r dr = K   2 qr 1 + q 2α2

(16.10)

c

Thus, by plotting [Is (q)]−1/2 against q 2 , one should obtain a straight line from which the ratio of the slope to the interception gives the αc2 . For a very dilute mixture where φ1  φ2 and φ1 → 1, αc equals to the average size of domains. According to Eq. (16.10), when q is very large, we have Is (q) ∼ q −4

(16.11)

Thus if q 4 I (q) is plotted against q 2 , the resulting profile should lever off with a constant value, which is referred as Porod’s law. In a two-phase system with diffusive interface, a deviation from Porod’s law could be observed because the actual correlation function is not presented as an exponential form. 16.2.2 16.2.2.1

Experiment Setup

A simple, small-angle light scattering (SALS) arrangement consisting of light source, optical align element, sample stage, screen and detector is shown in Figure 16.1. For SALS, a laser source is normally used, producing the uniform light wavelength and stable light power (within 5%). Usually He-Ne laser source with λ = 632.8 nm (red laser light) is involved. Recently, semiconductor laser (λ = 532 nm or 650 nm, green light) with an adjustable stable power output is also used as the light source for SALS. The power of light source used for SALS usually ranges from 1 mW to 50 mW. The optical elements involved normally include pin hole, lens and mirrors, which align the light path passing through the sample measured and project the scattering pattern onto the detector. The quality of the scattering image depends greatly on the characteristics and the alignment of the lenses. Sample films with thickness about 10 ∼ 100 μm are usually mounted between microscope slide cover slips and held on a hot stage equipped with a programmable temperature controller mounted on an optical bench of the SALS apparatus, which can provide rapid temperature control with the least temporal and spatial temperature variations. The acquirement of the scattering pattern is to convert the scattered light into a real image with a screen and transfer it to a detector with a high-quality camera lens. Photodiode array and charge-coupled device sensor CCD are commonly used as the detector to collect the scattering pattern. In the early days, the light intensity scattered by the sample was directly collected via a set of one-dimensional photodiode array. With the development of electrooptical technology, the photodiode array was replaced by the large-area CCD which could obtain the light intensity in two dimensions [6], in order to avoid the time-consuming alignment

Pin hole Polarizer

Laser

Figure 16.1

Analyzer

Sample

Detector

Schematic representation of experimental arrangement for light scattering.

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He-Ne Laser λ/4 plate Polarizing Beam-splitter Sample CCD Mirror

Figure 16.2 Plan view of the light-scattering apparatus. The light from the He-Ne laser impinged on the sample via an attenuator (quarter-wave plate and polarizing beam splitter) and glass plate oriented to extract, from the incident beam, an intensity-calibration beam. Reprinted from [6]. Copyright (1992) with permission from The American Physical Society.

process and lens aberration problems. In this case, the quality of screen is the vital factor in obtaining correct scattering data. The screen requires a nonreflecting surface with uniform and linear brightness response over a wide dynamic range. Different screens such as ground glass, copy paper and kraft paper may be used [7]. Figure 16.2 shows a small light scattering setup designed by Nakatani et al. [6]. It requires an objective lens with a large numerical aperture and a small focal length to collect scattered light, and a combination of lenses (e.g. two plano-convex lenses) to focus the image onto the detector. A beamstop is usually used to block the main beam, but light diffracted by the beamstop may distort the scattering pattern at small angles. These imperfections can severely limit the experimentally accessible range of angles over which the scattering intensity data are reliable. To lessen these defects, one needs to maximize the diameters of the lenses to reduce the spherical aberrations and carefully design and position the beamstop to reduce stray light. 16.2.2.2

Sample

The sample used for SALS could be prepared via solution casting, spin-coat and hot-compression. Some requirements must be obeyed when preparing a sample for SALS experiment. First, samples with low turbidity (high transmittance >90%) is preferred, i.e. thin samples with low concentration of minor phase and small refractive index difference. The thickness of the sample must be in a reasonable range in order to guarantee sufficient scattering intensity and prohibit the occurrence of multiple light scattering. Multiple light scatterings occur when a thick or concentrated sample is involved, which could be treated by the Mie scattering theory [8]. 16.2.2.3

Flow-light Scattering

The molecular structure and morphology of polymers produced in flow has now become a most interesting topic because of its great significance in determining the ultimate properties of polymer materials. Thus, the real-time evaluation of rheological behavior and morphology of multiphase polymer systems under flow is particular important both in academic research and industrial application [9]. During the past 20 years, many flow-light scattering (or rheo-light scattering) apparatus for investigating the influence of flow field on the phase behavior of polymer blends have been constructed based on commercial rheometers. The combination of rheology and light scattering techniques provides more constructive information in understanding the structure evolution of multiphase systems under conditions similar to the real processing procedure.

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Mirror CCD

Shear cell Screen

Oven

Laser

Figure 16.3 Schematic of a typical flow-light scattering apparatus. Reprinted from [10]. Copyright (1995) with permission from American Institute of Physics.

Figure 16.3 shows a typical schematic configuration of flow-light scattering apparatus used in many studies [10, 11]. The rheometer provided a well-defined flow field and recorded the data of the torque (shear stress) and the normal force. Various optically transparent rheometer geometries made of quartz such as parallel plates, cone and plate, and concentric cylinders were used. Although a parallel-plate geometry has the advantage of less optical distortion, the shear rate is not uniform along the radial direction. A cone and plate geometry provides a uniform shear rate to the sample, but the nonparallel surfaces of the geometry may distort the scattered image. Mirrors, reflection prisms, and quartz windows are usually used in order to allow the pass through of an incident laser light and the collection of scattering pattern. The Cartesian coordinate used in flow-light scattering experiments is given in Figure 16.4. The most complicated rheology-light scattering setups till now were probably designed by Han et al. [12] in 1996 and Hashimoto et al. [13] in 1999, respectively, in which the light scattering unit and microscopy unit were compactly integrated into a commercial rheometer and the different work modes in the machine could be switched readily, as shown in Figure 16.5. In such highly-integrated apparatus, the structure evolution of multiphase polymer systems obtained in a reciprocal space could be compared with that in the optical micrographs conveniently. Most flow-light scattering apparatus reported by different research groups z (Flow direction)

(a)

(b)

Laser beam

y y

x

x (c)

VV

HH P

Flow direction

HV

VH P

A

A P A

A

P

Figure 16.4 Cartesian coordinate used in flow-light scattering experiments. The propagation of the incident beam is along the shear gradient direction x axis, and the y axis and z axis is the vorticity direction and flow direction, respectively: (a) top view; (b) side view; (c) definition of four possible polarization conditions. Reprinted from [11]. Copyright (1995) with permission from Springer.

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A3 B17

C12

C11 C10 C8

C13

C9

C1

C14

B13

B12 B11 C7 B9 C6

B10 A2

A1

C2 C3

B16

B14 B15

B18

B3 A5

C4

B8

C5

B7 B6 B5 B4 B2

C

B

B1 D1

A4

A

Figure 16.5 Schematic diagram of rheo-optical SALS instrument designed by Hashimoto et al. [13]. Reprinted from [13]. Copyright (1999) with permission from American Institute of Physics.

were homemade with various structure designs. Recently, Alfa Corporation announced the production of a commercial rheo-light scattering setup based on their rheometer product. 16.2.3

Intensity Calibration

The goal of light scattering experiment is to obtain accurate relative scattering intensities over the entire accessible range of light scattering angles and at all azimuthal angles. The intensity calibration is crucial in obtaining the correct scattering data from experimental light scattering pattern, and cannot be overlooked and treated casually as experimental errors, which was found in some cases. The intensity corrections for the optics of the instrument include the effect of the refraction and the optical path of the light and the solid angle detected by each pixel of the CCD detector. When the incident light strikes a boundary between two media, the transmitted light intensity depends on the incident angle and the refractive indices of the materials. As a result, the value of the transmitted light intensity should be corrected or calibrated, especially for larger scattering angles. This correction involves solution of the Maxwell equations for the electromagnetic theory of light and is quite complex. Please refer to [14] for details. Here, only the angular calibration [15] is described briefly. Because of the refraction, the scattered light transmitted through the glass plate deviates from the actual scattering angle θ1 to the angle θ3 , as shown in Figure 16.6. The deviation of the scattered light depends on the thickness and the refractive index of the glass plate and the distance between the glass plate and the screen. For a specific scattering angle θ1 the distance d measured from the center of the beam was calculated using Snell’s law: d = d1 tan θ1 + d2 tan θ2 + d3 tan θ3

(16.12)

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d

Screen

θ1

d3 Glass Plate

θ3 θ2

d2 d1 Sample

Bottom Plate

Window

Laser

Figure 16.6 Schematic of the light path through the rheometer fixture and sample. Reprinted from [11]. Copyright (1995) with permission from Springer.

with θ2 = sin−1



   n1 n2 sin θ1 , θ3 = sin−1 sin θ2 n2 n3

(16.13)

where, n 1 , n 2 and n 3 are the refractive indices of the sample, glass plate, and air, respectively, and d1 , d2 and d3 , are the sample thickness, glass plate thickness, and the distance between the screen and the glass plate, respectively. The relationship between the pixel numbers and the scattering angle is determined by d and the magnification of the lens/detector system.

16.3 Phase Behavior of Multiphase Polymer Systems Studied by SALS 16.3.1

Thermodynamics

The phase behavior of polymer blends under equilibrium conditions has been extensively studied and different phase diagrams of polymer blends including LCST (Lower critical solution temperature) and UCST (Upper critical solution temperature) can be determined conveniently by SALS technique [16, 17]. When using SALS, the temperature under which the visible scattering phenomenon appeared is defined as the cloud point, which is reported to be close to the temperature at which binodal phase separation takes place. The accuracy of the cloud point is affected by the sample thickness, the heating/cooling rate and other factors [18]. The cloud point for a specific composition of polymer blend with LCST type would be drifted to a higher temperature at a rapid heating rate. The correct way to obtain the real cloud point is to measure the cloud point at various heating rates and then plot the cloud point versus the heating rate. By extrapolating the heat rate to zero, one could obtain the cloud point. For the convenience of comparison, the extrapolation may be ignored as long as the heating rate was kept constant. The SALS technique was utilized to probe the miscibility of homeopolymer blends or polymer solutions in the early time. In these studies, the miscibility originated from dipole–dipole interaction or specific interaction such as hydrogen bonding was discussed. The influence of temperature, component ratio, molecular architecture, molecular weight and distribution on the miscibility was discussed extensively [19, 20]. The results obtained from SALS experiments were usually compared with those predicted from various theories

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such as the Flory-Huggins theory [6, 21, 22] and equation of state (EOS) [16, 23]. The SALS technique has made a significant contribution to the theoretical development and the industrial application of polymer blends. Considerable emphasis was also placed on investigating how the nature, sequence and ratio of copolymer monomers play their roles in determining the miscibility of copolymer/homopolymer [24, 25] or copolymer/copolymer pairs [26, 27]. For example, by using SALS, Tang et al. [27] investigated the thermodynamic phase behavior of blends composed of two random copolymers of polystyrene-co-methyl methacrylate (SMMA) and polystyrene-co-acrylonitrile (SAN). The results indicated that SMMA and SAN are miscible within a certain copolymer composition ranges and exhibit LCST behavior. The segmental interaction energy parameters χS−MMA , χS−AN and χMMA−AN were calculated using a binary interaction model by fitting the lowest temperature in the calculated spinodal boundary from modified EOS theory to the light scattering data for several copolymer compositions. 16.3.2

Phase Separation Dynamics

Owing to its real-time and nondestructive features, the SALS technique has been utilized widely to probe the coarsening kinetics of phase-separation structures of polymer blends over the past decades [6, 28, 29]. 16.3.2.1

Spinodal Decomposition

Spinodal decomposition (SD) occurs when a binary mixture was deeply quenched into its instable region. In the early stages of SD, it was found that the time evolution of concentration fluctuations can be well described by the Cahn-Hilliard linearized theory [30]. In this stage, the Cahn-Hilliard theory predicts a single scattering maximum with constant position, along with an exponential growth of the scattered intensity I (q, t) with phase separation time t, qm (t) = qm (t = 0) I (q, t) = I (q, t = 0)e2R(q)/t

(16.14) (16.15)

where q is the wavenumber of the concentration fluctuations and qm is the predominant wavenumber for the fluctuations. The R(q) in Eq. (16.15) is the growth rate of the fluctuation of wavenumber q. For small q regime with q R0  1, where R0 is the individual polymer coil size, the Cahn-Hilliard theory permits estimation of the characteristic parameters describing the dynamics of the spinodal phase separation since   R(q)/q 2 = Dapp 1 − (q/qc )2

(16.16)

where Dapp is the apparent diffusion coefficient and qc is the maximum wavenumber of fluctuations that can grow. By plotting R(q)/q 2 as a function of q 2 , one can obtain Dapp from the intercept qc ,√from the value at which R(q)/q 2 = 0, and qm , the most probable wavenumber from the relation qm = qc / 2. Since the wavelength of the fluctuations, , corresponding approximately to two times the diameter of the phase domains, is related to the wavenumber q( = 2π/q), a minimum wavelength of fluctuations c that can grow, and a predominant wavelength m can be defined accordingly. Although the phase separation dynamics during SD can be accessed by other offline techniques such as digital image analysis [31], the SALS technique displays its superiority in monitoring the SD phase separation of polymer blends in real-time [32] and even at a molecular level [33]. Tang et al. [32] investigated the kinetics of SD phase separation for blends of the random copolymer poly(styrene-co-methyl methacrylate) (SMMA)

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and poly(styrene-co-acrylonitrile) (SAN) by using SALS. It was found that the time evolution of qm at the beginning of the late stages of SD phase separation followed the relationship qm ∝ t −1/3 , corresponding to an evaporation–condensation mechanism. Self-similar growth of the phase-separated structures at different timescales was also observed at the late stage. Wacharawichanant et al. [33] investigated the effects of low molar mass liquid crystals (LC) and lubricants on the molecular motion observed through early stage phase separation via spinodal decomposition in melt mixed PMMA/SAN blends by using light scattering techniques. It was found that the major effect of LC was to increase the molecular mobility of the blends and the Cahn–Hilliard apparent diffusion coefficient of the blend. Lubricants such as GMS and zinc stearate could also improve the mobility of the blend but to a lesser extent, and the effect did not increase at a higher concentration. It was proposed that the improved molecular mobilities might arise from changes in molecular segmental dynamics rather than thermodynamic effects. 16.3.2.2

Nucleation and Growth

Nucleation and growth occurred when a partial miscible polymer blend underwent a temperature jump into the metastable region of phase diagram. Dispersed phases were usually formed in this case. Using a time-resolved light scattering, Cumming et al. [29] found that the kinetics of domain growth for off-critical mixtures of polyisoprene and poly(ethylene-propylene) in metastable region is a heterogeneous nucleation process. It was found that the scattering curve of a mixture with ϕ = 0.48 after a shallow quench could be fitted by the scattering function from a near monodisperse spheres with Gaussian-distributed radii, as shown in Figure 16.7. The time dependence of the radius of the dispersed phases in the very earliest times obeyed approximately a r ∼ t 1/2 law. However, after a short time, the continuous phase surrounding the droplets developed into a morphological pattern reminiscent of that of spinodal decomposition (SD), which led to the formation of a characteristic scattering ring and a t 1/3 law in the domain growth similar to that of SD.

POLYDISPERSITY %

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INTENSITY (arb. units)

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10–2

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6

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Figure 16.7 I(q) for scattering from a nearly monodisperse distribution of spheres. The open circles are data and the curves are a model calculation from spheres with Gaussian-distributed radii, with polydispersity (−v/p) 0%, 3%, 7%, and 12%. Reprinted from [29]. Copyright (1990) with permission from The American Physical Society.

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104

double peak (761 min.) double peak (528 min.)

PB/PI =65/35 wt.%/wt/% T =33 °C

103

102

186 min. 95 min.

1

10

307 min.

(a)

40 min. before quench

100

Figure 16.8 Time change in the light scattering profiles of PB/PI(65/35) after temperature jump to 33◦ C. Reprinted from [34]. Copyright (1999) with permission from American Institute of Physics.

16.3.2.3

Percolation–cluster Transition

In critical mixtures, each phase keeps percolation even in the late stage of the phase separation where the interface between the separated domains is well defined. However, in off-critical mixtures even where their phase separation took place via SD, the minority phase cannot keep percolation any longer with the development of phase separation,, and the systems undergo a percolation-to-cluster transition (PCT) [12, 34, 35]. It was found that the percolation-to-cluster transition is influenced not only by the composition but also by the quench depth, because volume fraction of each phase also depends on the quench depth. Takeno et al. [34] found that two scattering peaks occurred when an off-critical PB/PI (65/35) mixture was quenched into the unstable region. The evolution of first peak wave number qm1 with the time followed the power law of percolated structures (qm1 ∼ t −a ; α = 0.9), while that of the second peak wave number qm2 was very slow, which corresponds to the growth of cluster of the droplets after PCT, as shown in Figure 16.8. Furthermore, the real space pictures and the light scattering analysis revealed that these clusters of small spherical droplets and the locally percolating domains could coexist for a very long time. Demyanchuk et al. [36] investigated the phase separation kinetics of an off-critical mixture of polystyrene and poly(methylphenylsiloxane). The results from the light scattering experiments were correlated with the images obtained by the optical microscopic observation in order to find characteristic features of the scattering intensity during the percolation-to-droplets morphology transition. It was found that when the large bicontinuous network starts to break up into disjointed elongated domains a second peak in the scattering intensity appears. Finally, both peaks merge into a single peak at zero wave vector, indicating a complete transformation of elongated domains into spherical droplets of varying sizes. 16.3.3

Reaction-induced Phase Separation

The reaction-induced phase-separation behavior and the structure development of thermoset/thermoplastic mixtures have been widely investigated on line by time-resolved small angle light scattering (TR-SALS) [37–44]. The morphology and final properties of the curing system were controlled by the competition

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(a) 170 °C

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qm /μm–1

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140 °C

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180 °C

160 °C 150 °C 140 °C

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1000

2000

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time /sec Figure 16.9 Temporal evolution of (a) qm and (b) Im for unfilled samples (hollow symbols) and filled samples (solid symbols) at cure temperatures of (, ) 140, (•, ◦) 150, (, ) 160, (, ) 170, (䉬, ) 180, and () 190◦ C. Reprinted from [49]. Copyright (2007) with permission from John Wiley & Sons.

between the curing reaction and the phase separation. It was verified that the phase separation of these systems followed the spinodal decomposition (SD) mechanism through analyzing the scattering pattern evolution and the co-continuous morphology observed by optical microscopy (OM). The curing temperatures and compositions of these epoxy systems were considered to be the determining factors for the domain size and the phase-separation behavior [42]. A four-stage model was discussed: (a) reaction prior to phase separation where the critical extent of reaction is the chemical quench for spinodal decomposition; (b) early stage spinodal decomposition where the scattering intensity increases exponentially with time and the CahnHilliard linear theory can be applied to the system; (c) late stage spinodal decomposition where the peak position shows a power law with time, and (d) apparent phase dissolution where the scattering intensity starts to decrease with time caused by a large change in the refractive index of the epoxy resin [44]. Both theoretical and experimental studies indicated that filler particles and fibers affect the morphology and phase-separation behavior of polymer blends, especially when the fillers are preferentially wettable to one component of the blends [45–48]. Zheng et al. [49] investigated the influence of nanofillers in the reactioninduced phase separation kinetics and the final morphology with TR-SALS. The phase-separation kinetics were analyzed by means of the temporal evolution of scattering vector qm and scattering intensity Im at the scattering peak. The organically modified layered silicates obviously facilitated an earlier onset of phase separation but reduced the phase-separation rate (Figure 16.9) and greatly retarded the domain coarsening process in the late stage of spinodal decomposition. The temporal evolution of both qm and Im followed the power law. Tang et al. [50] studied the cure-reaction-induced phase separation in attapulgites (ATTs)/epoxy/poly(ether sulfone) (PES) ternary hybrid nanocomposites using TR-SALS. The phase morphologies depended on the curing condition, the composition, the molecular weights and distribution of thermoplastics. ATT particles arranged at the interfaces blocked the movement of the interfaces, which made the coarsening process

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incomplete and inhibited the occurrence of the secondary phase separation. The incorporation of ATT particles did not affect the scope of the critical composition of the epoxy/PES blends but changed the final phase morphology which had a small phase domain and silklike PES-rich phase. 16.3.4

Phase Behavior of Polymer Blends Under Shear Flow

Following the intensive investigations of phase behavior of polymer blends under equilibrium conditions for many years [51, 52], the shear flow induced structure of polymer blends has rapidly become a significant subject of research because the shear flow which consisted of almost all industrial processing methods of polymer blends, such as injection molding and extrusion molding, would have a profound effect on the microstructure and ultimate properties of polymer blends [53–55]. In order to study the phase behaviors of polymer blends under shear flow, many characterization technique are used [56–58]. However, among them SALS is one of the most widely-used approaches, due to its nondestructive, well-established theory and technique for measuring in situ the change of the size, shape, degree of order and orientation of polymer species in complex fluids [59–62]. 16.3.4.1

Miscibility of Polymer Blends Under Shear Flow

The momentous influence of shear flow on the miscibility of partially miscible polymer blends includes the magnitude and the direction changes of the phase boundary. Shear-induced phase mixing (SIM), shearinduced demixing (SDM) and sometimes the co-existence of the SIM and SDM in the same blend were reported [63–67]. The SIM phenomenon is common in polymer blends [67, 68]. Kammer et al. [67] found that PMMA/SAN blend demonstrated a SIM behavior, as shown in Figure 16.10, that the shear flow suppresses phase separation and enlarges the homogeneous region of the blend. 240

2-phase

Temperature (°C)

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0 0.4 1.0 2.0 4.0

200

s–1 s–1 s–1 s–1 s–1

180 1-phase 100

0

20

40 60 SAN (wt%)

80

100

Figure 16.10 Cloud-point curves of PMMA/SAN blend under shear flow. Reprinted from [67]. Copyright (1991) with permission from Elsevier.

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80 1 phase

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0

20

40

60

80

100

Weight % PS

Figure 16.11 Elsevier Ltd.

The phase diagrams of PS/PVME blends under shear flow. Reprinted from [70]. Copyright (1999)

Soontaranun et al. [69] investigated the SIM behavior through the use of a generalized Gibbs energy of mixing modified by a stored energy term by the polymeric molecules during flow. It was indicated that the SIM behavior occurs when the excess stored energy is negative, but the phase separation could take place if the excess stored energy is positive. It was reported that shear flow could also reduce the homogeneous region of the polymer blends, and sometimes SDM and SIM may coexist in the same blend [70]. Figure 16.11 shows that the phase diagrams of PS/PVME blends were strongly influenced by the shear flow applied. At low shear rate (γ˙ = 0, 5, 8 s−1 ), the PS/PVME blends showed SDM behavior. However, when the shear rate was up to γ˙ = 14 s−1 , the homogeneous region reduced and the SIM occurred. Under certain conditions, polymer blends may show a quite complex miscibility behavior at shear flow [55, 59, 65, 66]. Lei [55] observed the effect of shear flow on the cloud point of the polymer blend PS/PVME. Figure 16.12 gives the phase separation temperature of PS/PVME with a critical composition of 70v% PVME as a function of shear rate. The results suggested that whether the shear started from low or high temperatures, there are two cloud point temperatures, which is similar to the experimental results of Fernandez [59]. At a higher shear rate (>8 s−1 ) such two-cloud-point-phenomenon disappeared and one phase separation temperature approximately 20 K higher than the quiescent cloud point was observed and it did not show any pronounced change with shear rate. Madbouly’s recent results of PS/PVME showed the same trend at higher shear rates [66].

16.3.4.2

The Dynamics Study of Polymer Blends Under Shear Flow

A significant difference between the dynamics of polymer blends at a quiescent state and those under the shear flow is that the shear flow could make the disperse phase orientated, which brings anisotropy to the polymer blends. To get the relevant orientation information of polymer blends under shear flow is very difficult by conventional characterization means. The SALS method could effectively record the changes of the light intensity from different directions (parallel and vertical to the flow direction), thus giving the orientation information about the phase structure easily. The level of the orientation of the disperse phases could be illustrated by the shape of the scattering patterns collected.

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o

T-Ts ( C)

20

10

o

0

110 C o 70 C

-10 0

10

20

30

40

-1

Shear rate (s )

Figure 16.12 The T-Ts vs. shear rate curve for PS/PVME with a critical composition of 70v% PVME at 383K (•) and 343K (). Ts is the quiescent cloud point. T-Ts>0 means shear-induced mixing. Reprinted from [55]. Copyright (2003) with permission from John Wiley & Sons.

For example, the light scattering from a phase-separated structure of PS/PVME blend (Figure 16.13(a)) under static state is usually an isotropic circular pattern (Figure 16.13(b)) [59]. Upon shearing the sample, the elongation of the dispersed phases in the flow direction [Figure 16.14(a)] resulted in a gradual transformation of the isotropic scattering pattern into an anisotropic oval shape and finally a streak-like pattern at a high shear rate normal to the flow direction. Figure 16.14(b) shows an oval-like scattering pattern in the intermediate stage of the transformation that was formed by a superimposition of an isotropic circular pattern with an oriented intense streak pattern. Nakatani et al. [63] investigated the influence of shear rate on the morphology of PS/PB(50/50)/dioctyl phthalate(DOP) 3.3 wt% under the steady flow. Figure 16.15(a) and (b) shows that near isotropic circular patterns were presented when the shear rate was very small, and these patterns would be developed into anisotropic patterns and finally changed to sharp streaks with the increasing of shear rates (Figure 16.15(c)–(e)). With the

(a)

( b)

Figure 16.13 (a) Optical micrograph of PS/PVME blend in quiescent conditions; (b) isotropic scattering signal obtained from a sample similar to that in (a). Reprinted from [59]. Copyright (1995) with permission from Elsevier.

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(b)

Figure 16.14 (a) Optical micrograph obtained for a sample of PS/PVME under shear flow; (b) two-dimensional scattering pattern for a sample similar to that in (a). Reprinted from [59]. Copyright (1995) with permission from Elsevier.

further increasing of shear rates, the streak-like patterns got less distinct and disappeared (Figure 16.15(i)), suggesting the shear-induced homogenization. The data obtained from SALS would also provide more information about the phase separation dynamics of polymer blends in different directions under shear flow [71]. Figure 16.16 is the relationship between ln [I (q) − Ibaseline ] and time, where q is scattering vector and I is the scattering intensity. From Figure 16.16(a) PS/PB(50/50)/DOP 3.3wi% ΔT(0)=8k x

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Figure 16.15 The SALS patterns of PS/PB(50/50)/DOP at various shear rates (s−1 ) under the steady flow. Reprinted from [63]. Copyright (1990) with permission from American Institute of Physics.

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Figure 16.16 ln[I(q) − Ibaseline ] versus time as a function of q. Reprinted from [71]. Copyright (1992) with permission from Elsevier.

and (b), it could be found that the evolution of the scattering intensity with the time is different in the direction normal to the shear flow from that in the direction parallel to one. Figure 16.17(a) showed that the size growth of the disperse phase is quite similar in both directions normal or parallel to the shear flow. It could be realized from Figure 16.17(b) that the most probable domain size of the disperse phase (corresponding to qm ) comes into being more quickly in the parallel direction than in the normal direction. 16.3.5

Multi-scale Approaches in Studying the Phase Behavior of Polymer Blends

It must be noted that the phase separation phenomena in various multicomponent systems usually generate microstructures with diverse dimensions ranging from 0.01 μm to 1 mm, which may exceed the length scale window of light scattering. For example, over-coarsened domains will lead to the statistic failure

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of light scattering, while the scattering signal produced by nano-scale microstructures is usually beyond the maximum measurable scattering angle of light scattering. Thus a suitable multi-scale characterization approach is required in some circumstances in order to probe the structure evolution of multiphase systems in an extended space window. Various techniques, such as atom force microscopy (AFM) [72] and in situ optical microscopy (OM) [73, 74] which may serve as important supplementary methods for light scattering, could provide a more complete morphological description in the real space. Actually, the light scattering pattern from a multiphase system corresponds directly to the optical Fourier transform of the microstructure in the real space [75]. Through the fast Fourier transforms (FFT) of micrographs, it is possible to investigate the phase behavior of multiphase systems by using the structural analysis methods similar to that of light scattering which have been discussed in previous sections. For example, Higgins et al. [72] investigated the bulk spinodal decomposition in thick film of tetramethyl bisphenol A polycarbonate (TMPC)/PS blend, which was shown to induce a surface roughness that was resolvable by AFM. The structure factors obtained from a 2D-FFT of ‘bulk’ AFM images were compared to those from 3D analysis of time-resolved light scattering. Han et al. [73] have investigated the relaxation of a binary polybutadiene/polyisoprene blend after the cessation of steady-state shear using both in situ optical microscopy and light scattering. The light scattering patterns were directly related with the real space morphology through fast Fourier transforms of the micrographs.

16.4 On-line Morphological Characterization of Polymer Blends Light scattering techniques have been proven to be very efficient for characterizing the morphology of polymer blends under shear flow [76, 77]. The main advantage of light scattering techniques – immediate response obtained during a shear flow without disturbing or rupturing the samples measured that could occur by using other conventional techniques like electron microscopy – has made it a preferable tool in monitoring on-line the morphology development of polymer blends during extrusion or mixing. An apparatus that utilized a commercial twin-screw extruding device to feed molten polymer into a narrow slit die was developed to directly probe the morphology of various polymer blend systems at the exit of the extruder by using SALS and optical microscopy (OM) together (Figure 16.18) [78–80]. This instrument was also used to investigate in situ the formation, orientation and morphology development of liquid crystalline polymer (LCP) fibers in a compatibilized poly(ethylene terephthlate) (PET) matrix during processing under different interfacial modification and processing conditions [81]. To follow the development of the polymer blend morphology along the extruder, a light scattering device was designed to mount on a twin-screw extruder at different locations along the barrel, as shown in Figure 16.19 [82, 83]. The light-scattering cell is directly connected to the extruder by means of special connectors. Hence, a small amount of molten polymer was able to deviate from the main flow inside the extruder to the lightscattering device. The deconvolution of the intensity profile gave access to the droplet size distribution and made it possible to study the influence of processing parameters such as feed throughput, RPM, screw profile and temperature. The difference in terms of light scattering intensity profile between reactive and nonreactive blends provided deeper insight into how the morphology is affected by the functionality (number of reactive functions per chain) and the processing parameters. For traditional SALS, the central speckle usually covers the scattering information with thin samples and the multiple scattering complicates the process of scattering information with thick samples. Back small angle laser scattering (BSALS) is the scattering part in reverse, which is independent of the central speckle; therefore it is thought to provide direct information on the formation, dissolution, and deformation of molecular aggregates [84]. The schematic diagram of the BSALS apparatus is shown in Figure 16.20. The mechanism of the formation and evolution of phase structure of binary blends of polypropylene (PP) with

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POLYMER

∧ Z ∧ X Y

Figure 16.18 The flow through the slit die is in the x direction, with the velocity gradient in the y direction and vorticity in the z direction. Sapphire optical windows in the top and bottom of the die permit optical monitoring during processing. The light scattering measures the projection of the structure factor onto the x–z plane. The focal plane of the microscope is parallel to this plane and can be adjusted as a function of y. Reprinted from [79]. Copyright (1998) with permission from John Wiley & Sons.

poly(cis-butadiene) rubber (PcBR) were investigated by a BSALS on-line system and on-line sampling in a mixer [62, 84, 85]. Structure parameters such as average chord length Lscatt and integral invariant Q were calculated to describe the relationship between the phase evolution and processing conditions. The velocity constant of the dispersed phase dimension variation k = dQ/dt was used at the early stage to discuss the relationship with different volume fractions of dispersed phase. The Lscatt of PP/PcBR blends in the early dispersion stage plotted against blending time in double-logarithmic scale could be linearly fitted, which meant the mechanism of phase crushing in PP/PcBR blends obeys the phase dispersion theory (brittle fracture theory). Therefore, this theory could be employed to study the kinetics of the formation and evolution of phase structures in PP/PcBR blends. Sheng et al. [86] also applied this online, nondestructive technique to investigate the phase morphology of polyethylene (PE)/polyamide 1010 (PA1010) immiscible blends along the extruder during twin-screw extrusion. The phase size was characterized by the correlation length ac , Lscatt , and average diameter of dispersed phase Dscatt according to the Debye theory. The dispersed particle size became progressively smaller along the screw length. The good agreement of BSALS with SEM showed that BSALS online system is a powerful method to investigate the phase size during twin-screw extrusion.

P2a P2b P3a

P3b P4a

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Figure 16.19 Screw profile of the twin screw extruder. P2a, P2b, P3a, P3b, P4a, P4b and P5a are the different possible locations of the light scattering device. The light scattering cell in mounted on P4b. Reprinted from [82]. Copyright (2002) with permission from John Wiley & Sons.

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Figure 16.20 Schematic representation of the backscattering apparatus: (L) He-Ne laser; (f) density filter; (P1, P2) Polaroid to change the polarization direction; (M) half-transparent reflective mirror; (S) sample; (O) optical system; (A) analysis system. Reprinted from [84]. Copyright (1997) with permission from Elsevier.

16.5 Light Scattering Characterization of Other Multiphase Polymer Systems 16.5.1

Gelation

The various physicochemical properties of gels are strongly correlated to the structure of crosslinking polymer chains and bridging sites. The study of the dynamic process of gel formation is an attractive research topic to clarify the gel structure and the chemistry and physics of gelation. The light scattering technique is one of the most suitable techniques to investigate a gelation process without disturbing the gelling system [87, 88]. Studies on the gelation kinetics and the structure development have been widely carried out. Structure inhomogeneities, the rather unique properties of gels, introduced during crosslinking processes, are particularly important as the inhomogeneities usually increase the turbidity, cause inhomogeneous response in swelling and shrinking, and affect the thermal response of temperature-sensitive polymer gels, etc. Moreover, the determination of the gelation threshold is a key issue to controlling the mechanical properties of a gelling system. 16.5.1.1

Inhomogeneities in Polymer Gels

In contrast to polymer solutions, most gels are inhomogeneous structures as the polymer chains in the network lose the freedom to travel freely in the phase space and are arrested in a limited phase space. The formation of heterogeneous structures of a crosslinked polymer gel has been the subject of great interest since the physical properties of gels such as permeability, elasticity, and optical and phase properties are directly affected by the structural inhomogeneities and extensive studies have been done in chemical gels using static light scattering and dynamic light scattering [89–101]. Okay et al. [93] observed a critical polymer network concentration existed where the degree of the inhomogeneity of poly(acryl amide) (PAAm) gels reached a maximum. Oppermann et al. [94] investigated the influence of crosslinker reactivity on the formation of inhomogeneities in PAAm hydrogels and related this to crosslinking efficiency. The effect of degree of crosslinking on spatial inhomogeneity in N-isopropylacrylamide-co-acrylic acid (NIPA/AAc) weakly charged copolymer gels [100] and partially charged PAA gels [92] was also investigated. By systematically researching the preparation conditions, Oppermann et al. revealed that spatial inhomogeneities grow stronger with rising preparation temperature (Tprep ) of the gels, increasing crosslinker concentration, and decreasing monomer concentration in poly(N-isopropylacrylamide) (PNIPA) hydrogels [95]. Based on the extensive investigation on the structure and dynamics of PAAm and PNIPA gels prepared at various conditions, Shibayama et al.

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Figure 16.21 Inhomogeneities in polymer gels. Reprinted from [102]. Copyright (2002) with permission from the Chemical Society of Japan.

[102, 103] described the diffusion coefficient and the scattered intensity component from the thermodynamic fluctuation by employing the partial heterodyne method. They proposed that there are several types of inhomogeneites in gels, as shown in Figure 16.21, which play significant roles to characterize gels and gelation thresholds. The spatial inhomogeneities are nonrandom spatial variations of crosslink density in a gel, which result in anomalous scattering. The topological inhomogeneities represent defects of network, which affect the dynamics and swelling behavior of gels. The connectivity inhomogeneities are dependent on cluster size, distribution, and architecture of polymer chains. The mobility inhomogeneities correspond to variations of local degree of mobility by introduction of crosslinks, which is the reason why scattering speckle appears exclusively in the gel state. Chemical gels are irreversible and normally formed by copolymerization of monomers with crosslinking agents. In contrast, in physical gels, polymer chains are interconnected via physical interactions, such as van der Waals, electrostatic attraction, and hydrogen bonding, to form reversible gel networks. It has been shown in thermoreversible physical gelling systems consisting of poly(vinyl alcohol) and Congo Red (PVA/CR) that the scattering speckles, i.e. random fluctuations in the scattered intensity with respect to sampling points, appeared exclusively in the gel state, i.e. T < T gel and CPVA > CPVA,gel . The appearance of speckles indicated that frozen inhomogeneities observed in chemically crosslinked gels also existed in thermoreversible physical gels [104, 105]. However, Wu et al. found that the static nonergodicity is not intrinsic for all thermally reversible physical gels by investigating the PNIPAM hybrid gels [106]. 16.5.1.2

Sol–gel Transition

Fang et al. [107] reported that dynamic light scattering is sensitive to the sol–gel transition of polymer solutions by chemical crosslinking. Martin et al. [87, 88] probed the sol–gel transition during the formation of silica gels by dynamic light scattering. They first reported on a power-law time decay of the intermediate scattering function, and compared with the data obtained by rheology test. Although many gelling systems exhibit a clear power-law behavior in the intensity correlation function near the sol–gel transition, power-law behavior is not observed in all cases as an indication for the gelation threshold [108]. It has been observed that the terminal relaxation time of concentration fluctuations does not diverge at the gel point, but only when the gel has become sufficiently dense [109, 110]. Shibayama et al. [102] used the fundamental analysis of the dynamic light scattering in nonergodic media [111] to investigate the gel formation by considering the loss of ergodicity during the gelation process. Nondestructive and real-time methodology using time-resolved dynamic light scattering (TR-DLS) has been employed to determine the gelation threshold, critical dynamic near gelation threshold, mechanism of gelation, and architecture of gelling cluster [102, 103]. Shibayama et al. found that light scattering intensity exhibited a characteristic rise at the gel point in various types of gelling systems, such as chemically gelling systems (NIPA gels and PAA gels), inorganic gels (TMOS),

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polymer hybrid (TMOS/DMAA), polymer-ion complexes (PVA/CR) undergoing reversible sol-gel transition, and natural polymers (gelatin and β-lactoglobulin) and bulk polymerization of styrene and divinylbenzene mixtures, etc. [112–115], where the polymer concentration is always above the chain-overlap concentration. Therefore, they proposed that the gel point is determined as a point at which one of the following features is observed: (1) an abrupt increase in scattering intensity; (2) a power-law behavior of the intensity–time correlation function (ITCF); (3) a characteristic broadening in the distribution function; and (4) suppression of the initial amplitude of ITCF. Recently, novel types of gels have been developed, such as slide-ring gel [116], nanocomposite gel [117], and double network gels [118]. Shibayama et al. [119, 120] investigated the gelation mechanism of PNIPAclay nanocomposite gels (NC gel) by dynamic light scattering and contrast variation small-angle neutron scattering (SANS). It was found that the gelation mechanism of NC gels is similar to that of conventional gels made with organic crosslinker (OR gels). 16.5.2

Crystallization

Light scattering is a particularly sensitive technique for detecting small density or orientation variations in initially homogeneous systems and therefore has been used in time-dependent studies of polymer blend crystallization by combining with OM, SEM, transmission electron microscopy (TEM), DSC, small-angle X-ray scattering (SAXS), wide-angle X-ray diffraction (WAXD), etc. It is well known that the morphology of polymer blends containing a semicrystalline and an amorphous, or two semicrystalline components is highly affected by the competition of the kinetic process of the liquid–liquid (L–L) phase separation and the crystallization of the system. When a polymer blend melt is quenched to a temperature below the spinodal temperature but above the melting temperature of the crystalline component, the morphology and structure of the polymer blend is mainly determined by the phase separation kinetics. While the temperature is quenched below both the spinodal point and the melting temperature of the crystalline component of the blend, crystallization may occur simultaneously and compete with L–L phase separation. Therefore, the final morphology of the blend is controlled by these two competitive processes. 16.5.2.1

Crystalline-amorphous Polymer Blends

Lee et al. [121–124] investigated the L–L phase separation and its effect on the crystallization behavior in polymer blends. Results confirmed that L–L phase separation via SD and the crystallization rate decreased with increasing L–L phase-separated time. In polypropylene (PP) /ethylene–propylene–rubber (EPR) system [122], memory of L–L phase separation remained even after the crystallization and the crystallization proceeded only in PP-rich phases. The rapid crystallization for short L–L phase-separated time could be ascribed to the elevation of chain mobility of PP by relatively higher amounts of EPR in PP-rich phases. PET/ poly(hydroxyl ether of bisphenol A) (phenoxy) blends are basically immiscible over a wide range of compositions. During melt processing, interchange reactions between the phenoxy hydroxyl and ester type carbonyl groups took place, which enhanced the miscibility of the blends. It was also found that L-L phase separation occurred via SD during annealing in PET/phenoxy blends. After the formation of the domain structure, the blend slowly underwent homogenization by the interaction between the two components. The crystallization rate of PET significantly depended on the composition change of the separated phases and the change of the sequence distribution in the polymer chains determined by the level of the interchange reactions between the two polymers [125]. Sakurai et al. [126] investigated the effect of phase-separated domains on the spherulite growth and the crystallization behavior of linear low density polyethylene (LLDPE) with a rubber polymer. The spherulites were 10–20 times larger than the phase-separated domains. The long period of crystalline lamellae was smaller than phase-separated domains. Therefore, it was supposed that the phase-separated

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structure did not influence on the fibril growth, which is the fundamental step of spherulites growth. Namely, the fibril growth though a channel of phase-separated crystalline domains enables the overall growth of spherulites. PET/poly(ether imide) (PEI) blends exhibited simultaneous liquid–liquid phase separation and crystallization over a wide range of temperature and composition. Nonlinear spherulite growths were observed for the blends at higher crystallization temperatures of 210◦ C and 220◦ C, while the growths were basically linear below 210◦ C. The nonlinear growth behavior was discussed based on the competition between spherulite growth and spinodal decomposition [127]. The nonlinear growth behavior was alternatively suggested to arise from the rearrangement of the spinodal domains associated with coarsening or from the continuous segregation of PEI content into the growth front, which gradually diluted the PET content as the growth front advanced. The spherulite growths in immiscible isotactic polypropylene (i-PP)/EPR blends over the modulated spinodal domains have been investigated by Hashimoto et al. [128, 129]. This system exhibited linear growth at a lower Tc of 140◦ C, but a nonlinear growth at a higher Tc of 145◦ C. The nonlinear growth behavior was suggested to relate to the long-range rearrangement of spinodal domains and the segregation of EPR into the growth front. Light-scattering analysis quantitatively revealed a competition of the crystallization (liquid–solid phase transition) and the phase dissolution (liquid–liquid phase transition) in terms of the crystallization rate (from HV scattering) and the apparent diffusion coefficient for dissolution (from VV scattering) in poly(ε-caprolactone) (PCL)/poly(styrene-co-acrylonitrile) (SAN) blend [130]. According to the features of the spinodal peak position and peak intensity, qm,Vv and Im,Vv , the phase separation in PCL/poly(ethylene glycol) oligomer (PEGo) was divided into four characteristic stages as indicated in Figure 16.22 [131]. The interplay of phase separation and crystallization has also been studied in isotactic polypropylene (iPP)/ethylene-propylene diene terpolymer (EPDM) [132], iPP/poly(butylene terephthalate) (PBT)/copolycarbonate [133], iPP/partially hydrogenated oligo(styrene-co-indene) (HIS) [134], PCL/poly(D,L-lactic acid) (PDLA) [135], polyamide-12 (PA-12)/poly (styrene-co-acrylonitrile) (SAN) [136].

16.5.2.2

Crystalline–crystalline Polymer Blends

Compared with crystalline-amorphous blend systems, the crystalline morphology development and crystallization behavior in binary crystalline–crystalline polymer blends are more complicated. Moreover, the possibility of the interplay between the phase separation and crystallization kinetics introduces additional complexity in the analysis. Therefore, very few studies have been reported on crystalline–crystalline polymer blends using light scattering [137–142]. The crystallization studies revealed that the high-density polyethylene (HDPE) and LLDPE formed strong cocrystalline mass when melt blended in a single screw extruder. The Avrami exponent indicated that the merging of individual nuclei of HDPE and LLDPE occurred to form cocrystallites. The rate of crystallization depended greatly on the number of crystallites and their interfacial boundary in contact with the amorphous phase [143]. Wang et al. [26, 144] investigated the kinetic interplay between crystallization and L–L phase separation on structural development in poly(ethylene-co-hexene) (PEH)/poly(ethylene-co-butene) (PEB). By controlling the relative quench depths for liquid–liquid phase separation and crystallization, the growth kinetics of the characteristic length scales of the simultaneous ordering processes showed a crossover from crystallization dominated to phase-separation dominated behavior. In the two-phase region above the melting temperature of polyethylene crystal, interconnected bicontinuous structural morphology was observed. It was found that amorphous polycarbonate (PC) crystallized in several minutes by blending poly(ethylene oxide) (PEO). The crystallization of PC was probably induced by the up-hill diffusion of the L–L phase separation via spinodal decomposition. Due to the competitive progress of the L–L phase separation and the crystallization, connected spherulites were obtained and the melting temperature decreased with increasing annealing temperature [145].

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References 1. P. Debye, Light scattering in solutions, J. Appl. Phys. 15: 338–342 (1944). 2. B. Chu, Laser Light Scattering: Basic Principles and Practice, Academic Press, London (1991). 3. R. S. Stein and M. B. Rhodes, Photographic light scattering by polyethylene films, J. Appl. Phys. 31: 1873–1884 (1960). 4. W. Brown, Dynamic Light Scattering Principles and Development. Clarendon Press, Oxford (1993). 5. W. Brown, Light Scattering: Principles and Development, Clarendon Press, USA, Oxford (1996). 6. A. Cumming, P. Wiltzius, F. S. Bates and J. H. Rosedale, Light-scattering experiments on phase-separation dynamics in binary fluid mixtures, Phys. Rev. A 45: 885–897 (1992). 7. V. Abetz, Artifacts in light scattering experiments using opaque scattering screens on the example of spinodal decomposition, Macromolecules 29: 4804–4805 (1996). 8. J. Kratohvil, G. Deˇzeli´c, M. Kerker and E. Matijevi´c, Calibration of light-scattering instruments: A critical survey, J. Polym. Sci. 57: 59–77 (1962). 9. C. C. Han, Y. Yao, R. Zhang and E. K. Hobbie, Effect of shear flow on multi-component polymer mixtures, Polymer 47: 3271–3286 (2006). 10. R. Wu, M. T. Shaw and R. A. Weiss, A rheo-light-scattering instrument for the study of the phase behavior of polymer blends under simple-shear flow, Rev. Sci. Instrum. 66: 2914–2921 (1995). 11. J. L¨auger and W. Gronski, A melt rheometer with integrated small angle light scattering, Rheol. Acta 34: 70–79 (1995). 12. S. Kim, J. W. Yu and C. C. Han, Shear light scattering photometer with optical microscope for the study of polymer blends, Rev. Sci. Instrum. 67: 3940–3947 (1996). 13. K. Matsuzaka and T. Hashimoto, A rheo-optical apparatus for simultaneous detection of rheology, small-angle light scattering, and optical microscopy under transient, oscillatory, and continuous shear flows, Rev. Sci. Instrum. 70: 2387–2397 (1999). 14. R. S. Stein and J. J. Keane, The scattering of light from thin polymer films. I. Experimental procedure, J. Polym. Sci. 17: 21–44 (1955). 15. R. Wu, M. T. Shaw and R. A. Weiss, A rheo-light-scattering instrument for the study of the phase behavior of polymer blends under simple-shear flow, Rev. Sci. Instrum. 66: 2914–2921 (1995). 16. T. Shiomi, F. Hamada, T. Nasako, K. Yoneda, K. Imai and A. Nakajima, Thermodynamics of polymer blends of poly(vinyl methyl ether) and polystyrene, Macromolecules 23: 229–233 (1990). 17. N. Posharnowa, A. Schneider, M. Wunsch, V. Kuleznew and B. A. Wolf, Polymer-polymer interaction parameters for homopolymers and copolymers from light scattering and phase separation experiments in a common solvent, J. Chem. Phys. 115: 9536–9546 (2001). 18. M. Zuo, M. Peng and Q. Zheng, Study on nonlinear phase-separation for PMMA/a-MSAN blends by dynamic rheological and small angle light scattering measurements, J. Polym. Sci. Pol. Phys 44: 1547–1555 (2006). 19. D. R. Paul and C. B. Bucknall, Polymer Blends, Wiley, New York (2000). 20. F. Bates, Polymer-polymer phase behavior, Science 251: 898–905 (1991). 21. T. Nose, Kinetics of phase separation in polymer mixtures, Phase. Transit. 8: 245–260 (1987). 22. M. Gehlsen, J. Rosedale, F. Bates, G. Wignall, L. Hansen and K. Almdal, Molecular weight scaling in critical polymer mixtures, Phys. Rev. Lett. 68: 2452–2455 (1992). 23. R. Roe and W. Zin, Determination of the Polymer-Polymer Interaction Parameter for the Polystyrene-Polybutadiene Pair, Macromolecules 13: 1221–1228 (1980). 24. Q. Yang, Y. M. Mao, G. X. Li, Y. J. Huang, P. Tang and C. H. Lei, Study on the UCST behavior of polystyrene/poly(styrene-co-acrylonitrile) blend, Mater. Lett. 58: 3939–3944 (2004). 25. K. Yang, Q. Yang, G. X. Li, Y. J. Sun and Y. M. Mao, Phase behavior of blends of poly (vinyl methyl ether) with styrene-acrylonitrile in film, Mater. Lett. 59: 2680–2684 (2005). 26. H. Wang, K. Shimizu, E. K. Hobbie, Z.-G. Wang, J. C. Meredith, A. Karim, E. J. Amis, B. S. Hsiao, E. T. Hsieh and C. C. Han, Phase diagram of a nearly isorefractive polyolefin blend, Macromolecules 35: 1072–1078 (2002). 27. P. Tang, G. X. Li, J. S. Higgins, V. Arrighi and J. T. Cabral, Phase behaviour of SMMA and SAN blends using Flory’s equation of state theory, Polymer 43: 6661–6667 (2002).

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28. T. Hashimoto, J. Kumaki and H. Kawai, Time-resolved light scattering studies on kinetics of phase separation and phase dissolution of polymer blends. 1. Kinetics of phase separation of a binary mixture of polystyrene and poly (vinyl methyl ether), Macromolecules 16: 641–648 (1983). 29. A. Cumming, P. Wiltzius and F. S. Bates, Nucleation and growth of monodisperse droplets in a binary-fluid system, Phys. Rev. Lett. 65: 863–866 (1990). 30. C. Han and A. Akcasu, Phase Decomposition in polymers, Annual Reviews in Physical Chemistry 43: 61–90 (1992). 31. H. Tanaka and T. Nishi, Direct determination of the probability distribution function of concentration in polymer mixtures undergoing phase separation, Phys. Rev. Lett. 59: 692–695 (1987). 32. P. Tang, V. Arrighi, J. S. Higgins and G. X. Li, Kinetics of phase separation of poly(styrene-co-methyl methacrylate) and poly(styrene-co-acrylonitrile) blends, Polym. Int. 53: 1689–1692 (2004). 33. S. Wacharawichanant, S. Thongyai, S. Tanodekaew, J. S. Higgins and N. Clarke, Spinodal decomposition as a probe to measure the effects on molecular motion in poly(styrene-co-acrylonitrile) and poly(methyl methacrylate) blends after mixing with a low molar mass liquid crystal or commercial lubricant, Polymer 45: 2201–2209 (2004). 34. H. Takeno, E. Nakamura and T. Hashimoto, Heterogeneous percolation-to-cluster transition in phase separation of an off-critical polymer mixture, J. Chem. Phys 110: 3612–3620 (1999). 35. H. Takeno, M. Iwata, M. Takenaka and T. Hashimoto, Combined light scattering and laser scanning confocal microscopy studies of a polymer mixture involving a percolation-to-cluster transition, Macromolecules 33: 9657–9665 (2000). 36. I. Demyanchuk, S. A. Wieczorek and R. Holyst, Percolation-to-droplets transition during spinodal decomposition in polymer blends, morphology analysis, J. Chem. Phys 121: 1141–1147 (2004). 37. K. Yamanaka and T. Inoue, Structure development in epoxy resin modified with poly (ether sulphone), Polymer 30: 662–667 (1989). 38. B. S. Kim, T. Chiba and T. Inoue, Phase-separation and apparent phase dissolution during cure process of thermoset thermoplastic blend, Polymer 36: 67–71 (1995). 39. T. Inoue, Reaction-induced phase decomposition in polymer blends, Prog. Polym. Sci. 20: 119–153 (1995). 40. B. S. Kim, G. I. Nakamura and T. Inoue, Structure development in the phenolic resin poly(methyl methacrylate-coethyl acrylate) copolymer blends, J. Appl. Polym. Sci. 68: 1829–1837 (1998). 41. J. Xu, M. Holst, M. R¨ullmann, M. Wenzel and I. Alig, Reaction-induced Phase Separation in a Polysulfone-modified Epoxy-anhydride Thermoset, J. Macromol. Sci. Phys. 46: 155–181 (2007). 42. G. Poel, S. Goossens, B. Goderis and G. Groeninckx, Reaction induced phase separation in semicrystalline thermoplastic/epoxy resin blends, Polymer 46: 10758–10771 (2005). 43. Q. Zheng, K. Tan, M. Peng and Y. Pan, Study on the phase separation of thermoplastic-modified epoxy systems by time-resolved small-angle laser light scattering, J. Appl. Polym. Sci. 85: 950–956 (2002). 44. Y. Ishii and A. Ryan, Processing of poly (2, 6-dimethyl-1, 4-phenylene ether) with epoxy resin. 1. Reaction-induced phase separation, Macromolecules 33: 158–166 (2000). 45. H. Tanaka, A. J. Lovinger and D. D. Davis, Pattern evolution caused by dynamic coupling between wetting and phase separation in binary liquid mixture containing glass particles, Phys. Rev. Lett. 72: 2581–2584 (1994). 46. V. V. Ginzburg, F. Qiu, M. Paniconi, G. Peng, D. Jasnow and A. C. Balazs, Simulation of hard particles in a phase-separating binary mixture, Phys. Rev. Lett. 82: 4026–4029 (1999). 47. G. Peng, F. Qiu, V. V. Ginzburg, D. Jasnow and A. C. Balazs, Forming supramolecular networks from nanoscale rods in binary, phase-separating mixtures, Science 288: 1802–1804 (2000). 48. F. Qiu, G. Peng, V. V. Ginzburg, A. C. Balazs, H. Y. Chen and D. Jasnow, Spinodal decomposition of a binary fluid with fixed impurities, J. Chem. Phys. 115: 3779–3784 (2001). 49. M. Peng, D. Li, Y. Chen and Q. Zheng, Effect of an organoclay on the reaction-induced phase-separation kinetics and morphology of a poly(ether imide)/epoxy mixture, J. Appl. Polym. Sci. 104: 1205–1214 (2007). 50. L. Zhao, G. Zhan, Y. Yu, X. Tang and S. Li, Influence of attapulgites on cure-reaction-induced phase separation in epoxy/poly (ether sulfone) blends, J. Appl. Polym. Sci. 108: 953–959 (2008). 51. W. Guo and J. S. Higgins, Miscibility and kinetics of phase separation in polymer blends of tetramethyl-bisphenol-A polycarbonate and polystyrene, Polymer 31: 699–706 (1990).

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52. L. A. Utraki, Polymer Alloys and Blends, Carl Hanser, New York (1989). 53. K. Mazich and S. Carr, The effect of flow on the miscibility of a polymer blend, J. Appl. Phys. 54: 5511–5514 (1983). 54. S. Madbouly and T. Ougizawa, Phase behavior and morphology of poly (methyl methacrylate)/poly (α-methyl styrene-co-acrylonitrile) blends under quiescent and shear flow, Polymer Reviews 45: 19–58 (2005). 55. C. H. Lei, G. X. Li, Q. Yang, D. H. Chen and S. J. Jiang, Effect of shear flow on the phase-separation behavior in a blend of polystyrene and polyvinyl methyl ether, J. Polym. Sci. Pol. Phys 41: 661–669 (2003). 56. F. Larbi, M. Malone, H. Winter, J. Halary, M. Leviet and L. Monnerie, Fluorescence evidence for shear flow-induced miscibility in a blend of polystyrene and poly (vinyl methyl ether), Macromolecules 21: 3532–3534 (1988). 57. P. Keates, G. Mitchell, E. Peuvrel-Disdier and P. Navard, In situ X-ray scattering study of anisotropic solutions of hydroxypropylcellulose subjected to shear flow, Polymer 34: 1316–1319 (1993). 58. N. J. J. Aelmans, V. M. C. Reid and J. S. Higgins, Change of phase behaviour of SMA/PMMA blends during processing at high deformation rates, Polymer 40: 5051–5062 (1999). 59. M. L. Fernandez, J. S. Higgins and S. M. Richardson, Flow instabilities in polymer blends under shear, Polymer 36: 931–939 (1995). 60. W. B. Ouyang, G. X. Li and Q. Yang, Shear-induced apparent shift of phase boundary of binary polymer blends, Macromol. Theor. Simul 13: 178–188 (2004). 61. S. Madbouly, T. Ougizawa and T. Inoue, Effect of shear memory on the coarsening behaviour of polystyrene and poly (vinyl methyl ether) blend, Polymer 45: 4483–4489 (2004). 62. Y. Yang, Z. Xiao, X. Jiang and J. Sheng, Light scattering studies on phase behavior of PP/PcBR blends during melt-mixing, J. Macromol. Sci. Phys. 45: 1083–1098 (2006). 63. A. Nakatani, H. Kim, Y. Takahashi, Y. Matsushita, A. Takano, B. Bauer and C. Han, Shear stabilization of critical fluctuations in bulk polymer blends studied by small angle neutron scattering, J. Chem. Phys 93: 795–810 (1990). 64. I. Hindawi, J. Higgins and R. Weiss, Flow-induced mixing and demixing in polymer blends, Polymer 33: 2522–2529 (1992). 65. M. Fernandez, J. Higgins, R. Horst and B. Wolf, Complex miscibility behaviour for polymer blends in flow, Polymer 36: 149–154 (1995). 66. S. A. Madbouly, T. Ougizawa and T. Inoue, Phase behavior under shear flow in PMMA/SAN blends: Effects of molecular weight and viscosity, Macromolecules 32: 5631–5636 (1999). 67. H. Kammer, C. Kummerloewe, J. Kressler and J. Melior, Shear-induced phase vhanges in polymer blends, Polymer 32: 1488–1492 (1991). 68. J. W. Yu, J. F. Douglas, E. K. Hobbie, S. Kim and C. C. Han, Shear-induced ‘homogenization’ of a diluted polymer blend, Phys. Rev. Lett. 78: 2664 (1997). 69. W. Soontaranun, J. S. Higgins and T. D. Papathanasiou, Shear flow and the phase behaviour of polymer blends, Fluid. Phase. Equilibr 121: 273–292 (1996). 70. S. Madbouly, M. Ohmomo, T. Ougizawa and T. Inoue, Effect of the shear flow on the phase behaviour of polystyrene/poly(vinyl methyl ether) blend, Polymer 40: 1465–1472 (1999). 71. D. Waldow, A. Nakatani and C. Han, Anisotropic phase separation kinetics in a polymer blend solution following cessation of shear studied by light scattering, Polymer 33: 4635–4638 (1992). 72. J. T. Cabral, J. S. Higgins, T. C. B. McLeish, S. Strausser and S. N. Magonov, Bulk spinodal decomposition studied by atomic force microscopy and light scattering, Macromolecules 34: 3748–3756 (2001). 73. L. Kielhorn, R. H. Colby and C. C. Han, Relaxation behavior of polymer blends after the cessation of shear, Macromolecules 33: 2486–2496 (2000). 74. H. S. Jeon and E. K. Hobbie, Shear viscosity of phase-separating polymer blends with viscous asymmetry, Phys. Rev. E 63: 061403–061401 (2001). 75. H. Tanaka, T. Hayashi and T. Nishi, Application of digital image analysis to pattern formation in polymer systems, J. Appl. Phys. 59: 3627 (1986). 76. N. Remediakis, R. Weiss and M. Shaw, Phase structure changes in a sheared blend of high molecular weight polybutadiene and polyisoprene elastomers, Rubber Chem. Technol. 70: 71–89 (1996). 77. S. Kim, E. K. Hobbie, J.-W. Yu and C. C. Han, Droplet breakup and shear-induced mixing in critical polymer blends, Macromolecules 30: 8245–8253 (1997).

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78. S. Li, K. B. Migler, E. K. Hobbie, H. Kramer, C. C. Han and E. J. Amis, Light-scattering photometer with optical microscope for the in-line study of polymer extrusion, J. Polym. Sci. Pol. Phys 35: 2935–2943 (1997). 79. E. K. Hobbie, K. B. Migler, C. C. Han and E. J. Amis, Light scattering and optical microscopy as in-line probes of polymer blend extrusion, Adv. Polym. Tech. 17: 307–316 (1998). 80. K. B. Migler, E. K. Hobbie and F. Qiao, In line study of droplet deformation in polymer blends in channel flow, Polym. Eng. Sci. 39: 2282–2291 (1999). 81. F. Qiao, K. B. Migler, D. L. Hunston and C. C. Han, Reactive compatibilization and in-line morphological analysis of blends of PET and a thermotropic liquid crystalline polymer, Polym. Eng. Sci. 41: 77–85 (2001). 82. G. Schlatter, C. Serra, M. Bouquey, R. Muller and J. Terrisse, Online light scattering measurements: A method to assess morphology development of polymer blends in a twin-screw extruder, Polym. Eng. Sci. 42: 1965–1975 (2002). 83. C. Serra, G. Schlatter, M. Bouquey, R. Muller and J. Terrisse, Online light scattering measurements as a means to assess iInfluence of extrusion parameters on non-reactive polymer blend morphology: Experimental procedure and preliminary results, Can. J. Chem. Eng. 80: 1036–1043 (2002). 84. J. Zhou and J. Sheng, Small-sngle light backscattering of polymer blends. 1. Multiple-scattering, Polymer 38: 3727–3731 (1997). 85. Y. Yang, W. Li, Z. Xiao, Y. Li, X. Jiang and J. Sheng, Study on the mechanism of the formation and evolution of phase structure in PP/PcBR blends, J. Macromol. Sci. Phys. 45: 1039–1052 (2006). 86. Y. Chen, J. Sheng, N. Shen and H. Jia, Phase-size analysis for multiphase polymer systems during twin-screw extrusion by small angle light scattering online measurements-polyethylene/polyamide immiscible blends, Acta Polym. Sin. 2: 282–287 (2004). 87. J. Martin and J. Wilcoxon, Critical dynamics of the sol-gel transition, Phys. Rev. Lett. 61: 373–376 (1988). 88. J. Martin, J. Wilcoxon and J. Odinek, Decay of density fluctuations in gels, Phys. Rev. A 43: 858–872 (1991). 89. S. Mallam, F. Horkay, A. Hecht and E. Geissler, Scattering and swelling properties of inhomogeneous polyacrylamide gels, Macromolecules 22: 3356–3361 (1989). 90. J. Joosten, J. McCarthy and P. Pusey, Dynamic and static light scattering by aqueous polyacrylamide gels, Macromolecules 24: 6690–6699 (1991). 91. E. Matsuo, M. Orkisz, S. Sun, Y. Li and T. Tanaka, Origin of structural inhomogeneities in polymer gels, Macromolecules 27: 6791–6796 (1994). 92. A. Moussaid, S. Candau and J. Joosten, Structural and dynamic properties of partially charged poly (acrylic acid) gels: Nonergodicity and inhomogeneities, Macromolecules 27: 2102–2110 (1994). 93. M. Kizilay and O. Okay, Effect of initial monomer concentration on spatial inhomogeneity in poly (acrylamide) gels, Macromolecules 36: 6856–6862 (2003). 94. B. Lindemann, U. Shroeder and W. Oppermann, influence of the cross-linker reactivity on the formation of inhomogeneities in hydrogels, Macromolecules 30: 4073–4077 (1997). 95. J. Nie, B. Du and W. Oppermann, Influence of formation conditions on spatial inhomogeneities in poly (Nisopropylacrylamide) hydrogels, Macromolecules 37: 6558–6564 (2004). 96. R. Liu and W. Oppermann, Spatial inhomogeneities of polystrene gels prepared from semidilute solutions, Macromolecules 39: 4159–4167 (2006). 97. M. Shibayama, Y. Fujikawa and S. Nomura, Dynamic light scattering study of poly (N-isopropylacrylamide-coacrylic acid) gels, Macromolecules 29: 6535–6540 (1996). 98. F. Ikkai and M. Shibayama, Anomalous cross-link density dependence of scattering from charged gels, Phys. Rev. E 56: 51–54 (1997). 99. M. Shibayama, F. Ikkai, Y. Shiwa and Y. Rabin, Effect of degree of cross-linking on spatial inhomogeneity in charged gels. I. Theoretical predictions and light scattering study, J. Chem. Phys 107: 5227–5235 (1997). 100. F. Ikkai, O. Iritani, M. Shibayama and C. C. Han, Effect of degree of cross-linking on spatial inhomogeneity in charged gels. 3. Ionization effect, Macromolecules 31: 8526–8530 (1998). 101. F. Ikkai and M. Shibayama, Inhomogeneity control in polymer gels, J. Polym. Sci. Pol. Phys 43: 617–628 (2005). 102. M. Shibayama and T. Norisuye, Gel formation analyses by dynamic light scattering, B. Chem. Soc. Jpn. 75: 641–659 (2002).

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103. M. Shibayama, Universality and specificity of polymer gels viewed by scattering methods, B. Chem. Soc. Jpn. 79: 1799–1819 (2006). 104. F. Ikkai and M. Shibayama, Static inhomogeneities in thermoreversible physical gels, Phys. Rev. Lett. 82: 4946–4949 (1999). 105. M. Shibayama, M. Tsujimoto and F. Ikkai, Static inhomogeneities in physical gels: Comparison of temperatureinduced and concentration-induced sol-gel transition, Macromolecules 33: 7868–7876 (2000). 106. Y. Zhao and C. Wu, A hybrid polymer gel and its static nonergodicity, Macromol. Symp. 207: 37–45 (2004). 107. L. Fang, W. Brown and C. Konak, Dynamic light scattering study of the sol-gel transition, Macromolecules 24: 6839–6842 (1991). 108. M. Takeda, T. Norisuye and M. Shibayama, Critical dynamics of cross-linked polymer chains near the gelation threshold, Macromolecules 33: 2909–2915 (2000). 109. T. Nicolai and S. Cocard, Dynamic light-scattering study of aggregating and gelling colloidal disks, J. Colloid Interf. Sci 244: 51–57 (2001). 110. V. Lesturgeon, T. Nicolai and D. Durand, Dynamic and static light scattering study of the formation of cross-linked PMMA gels, The European Physical Journal B 9: 71–82 (1999). 111. P. Pusey and W. van Megen, Dynamic light scattering by non-ergodic media, Physica. A 157: 705–741 (1989). 112. T. Norisuye, M. Shibayama and S. Nomura, Time-resolved light scattering study on the gelation process of poly(Nisopropyl acrylamide), Polymer 39: 2769–2775 (1998). 113. T. Norisuye, M. Inoue, M. Shibayama, R. Tamaki and Y. Chujo, Time-resolved dynamic light scattering study on the dynamics of silica gels during gelation process, Macromolecules 33: 900–905 (2000). 114. M. Okamoto, T. Norisuye and M. Shibayama, Time-resolved dynamic light scattering study on gelation and gelmelting processes of gelatin gels, Macromolecules 34: 8496–8502 (2001). 115. M. Shibayama, S. Ozeki and T. Norisuye, Real-time dynamic light scattering on gelation and vitrification, Polymer 46: 2381–2388 (2005). 116. Y. Okumura and K. Ito, The polyrotaxane gel: A topological gel by figure-of-eight cross-links, Adv. Mater. 13: 485–487 (2001). 117. K. Haraguchi and T. Takehisa, Nanocomposite hydrogels: A unique organicinorganic network structure with extraordinary mechanical and swelling/deswelling properties, Adv. Mater. 14: 1120–1124 (2002). 118. A. Bafna, G. Beaucage, F. Mirabella and S. Mehta, 3D Hierarchical orientation in polymer-clay nanocomposite films, Polymer 44: 1103–1115 (2003). 119. M. Shibayama, J. Suda, T. Karino, S. Okabe, T. Takehisa and K. Haraguchi, Structure and dynamics of poly (N-isopropylacrylamide)-clay nanocomposite gels, Macromolecules 37: 9606–9612 (2004). 120. S. Miyazaki, H. Endo, T. Karino, K. Haraguchi and M. Shibayama, Gelation mechanism of poly (Nisopropylacrylamide)-clay nanocomposite gels, Macromolecules 40: 4287–4295 (2007). 121. J. Lee, W. Choi, Y. Kwon and K. Lee, Liquid–liquid phase separation and crystallization behavior of poly (ethylene terephthalate)/poly (ether imide) blend, Polymer 43: 2827–2833 (2002). 122. S. W. Lim, K. H. Lee and C. H. Lee, Liquid-liquid phase separation and its effect on crystallization in the extruded polypropylene/ethylene-propylene-rubber blend, Polymer 40: 2837–2844 (1999). 123. J. Lee, K. Lee and C. Lee, Liquid phase behavior and its effect on crystalline morphology in the extruded poly (ethylene terephthalate)/poly (ethylene-2, 6-naphthalate) blend, Polymer 43: 1897–1902 (2002). 124. J. K. Lee, J. E. Im, J. H. Park, H. Y. Won and K. H. Lee, Morphology development and crystallization behavior of a poly(ethylene terephthalate)/polycarbonate blend, J. Appl. Polym. Sci. 99: 2220–2225 (2006). 125. G. Kim, J. Lee, M. Choi and K. Lee, Morphology development and crystallization behavior of poly (ethylene terephthalate)/phenoxy blend, J. Polym. Sci. Pol. Chem 46: 223–232 (2007). 126. K. Hayasi, Y. Ejima, T. Shimizu, S. Kizaki, S. Hara, K. Yamamoto, S. Okamoto and S. Sakurai, Crystallization behavior of linear low density polyethylene in its blend with a rubber polymer as revealed by synchrotron SAXS/WAXS/Hv-SALS simultaneous measurements, Nihon Reoroji Gakkaishi 32: 179–187 (2004). 127. H. Chen, J. Hwang, J. Yang and R. Wang, Simultaneous liquid-liquid demixing and crystallization and its effect on the spherulite growth in poly (ethylene terephthalate) poly (ether imide) blends, Polymer 39: 6983–6989 (1998).

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128. N. Inaba, K. Sato, S. Suzuki and T. Hashimoto, Morphology control of binary polymer mixtures by spinodal decomposition and crystallization. 1. Principle of method and preliminary results on PP/EPR, Macromolecules 19: 1690–1695 (1986). 129. N. Inaba, T. Yamada, S. Suzuki and T. Hashimoto, Morphology control of binary polymer mixtures by spinodal decomposition and crystallization. 2. Further studies on PP/EPR, Macromolecules 21: 407–414 (1988). 130. P. Svoboda, D. Svobodova, T. Chiba and T. Inoue, Competition of phase dissolution and crystallization in poly(εcaprolactone)/poly(styrene-co-acrylonitrile) blend, Eur. Polym. J. 44: 329–341 (2008). 131. W. Chuang, U. Jeng, P. Hong, H. Sheu, Y. Lai and K. Shih, Dynamic interplay between phase separation and crystallization in a poly (?-caprolactone)/poly (ethylene glycol) oligomer blend, Polymer 48: 2919–2927 (2007). 132. C. Y. Chen, W. Yunus, H. W. Chiu and T. Kyu, Phase separation behaviour in blends of isotactic polypropylene and ethylene-propylene diene terpolymer, Polymer 38: 4433–4438 (1997). 133. M. Okamoto and T. Inoue, Crystallization kinetics in poly(butylene terephthalate)/copolycarbonate blend, Polymer 36: 2739–2744 (1995). 134. C. Lee, Non-linear spherulite growth and phase change at the spherulite growth front in isotactic polypropylene/partially hydrogenated oligo (styrene-co-indene) blend, Polymer 39: 5197–5201 (1998). 135. Y. Akpalu and Y. Lin, Multivariable structural characterization of semicrystalline polymer blends by small-angle light scattering, J. Polym. Sci. Pol. Phys 40: 2714–2727 (2002). 136. Y. J. Sun, Y. J. Huang, Q. Yang and G. X. Li, Thermal behavior of polyamide-12 and poly (styrene-co-acrylonitrile) blend, J. Macromol. Sci. Phys. 45: 1015–1023 (2006). 137. J. Penning and R. St. John Manley, Miscible blends of two crystalline polymers. 1. Phase behavior and miscibility in blends of poly (vinylidene fluoride) and poly (1, 4-butylene adipate), Macromolecules 29: 77–83 (1996). 138. K. Fujita, T. Kyu and R. S. J. Manley, Miscible blends of two crystalline polymers .3. Liquid-liquid phase separation in blends of poly(vinylidene fluoride) poly(butylene adipate), Macromolecules 29: 91–96 (1996). 139. Y. Okabe, T. Kyu, H. Saito and T. Inoue, Spiral crystal growth in blends of poly(vinylidene fluoride) and poly(vinyl acetate), Macromolecules 31: 5823–5829 (1998). 140. I. Isayeva, T. Kyu and R. S. Manley, Phase transitions, structure evolution, and mechanical properties of blends of two crystalline polymers: poly(vinylidene fluoride) and poly(butylene adipate), Polymer 39: 4599–4608 (1998). 141. K. J. Kim and T. Kyu, Miscibility, phase behavior, and Curie transition in blends of vinylidene fluoride/trifluoroethylene copolymer and poly(1,4-butylene adipate), Polymer 40: 6125–6134 (1999). 142. J. Lee, A. Prabu, K. Kim and C. Park, Phase separation and crystallization behavior of poly (vinylidene fluoride)/poly (1, 4-butylene adipate) blends under an electric field, Macromolecules 41: 3598–3604 (2008). 143. S. Rana, Crystallization of high-density polyethylene-linear low-density polyethylene blend, J. Appl. Polym. Sci. 69: 2599–2607 (1998). 144. H. Wang, K. Shimizu, H. Kim, E. K. Hobbie, Z.-G. Wang and C. C. Han, Competing growth kinetics in simultaneously crystallizing and phase-separating polymer blends, J. Chem. Phys. 116: 7311–7315 (2002). 145. M. Tsuburaya and H. Saito, Crystallization of polycarbonate induced by spinodal decomposition in polymer blends, Polymer 45: 1027–1032 (2004).

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17 X-ray Scattering Studies on Multiphasic Polymer Systems Z. Z. Denchev and J. C Viana IPC – Institute for Polymers and Composites, University of Minho, Portugal

17.1 Introduction At present, multiphase and multicomponent materials such as polymer alloys, blends, and composites consume over 80 wt% of all commercially-produced resins. The increase of this particular segment of the plastics industry is about three times faster than of the industry as a whole. The reason is that the modification by blending may improve significantly the resins’ mechanical performance and processability, being at the same time cost effective. Those few resins that are used without adding of other components are most frequently semicrystalline, i.e. they already have a multiphase structure that makes other modification less urgent [1]. Recently, a large window has opened for new structural applications of the multicomponent polymer systems with the advent of nanoscale filled polymer composites. Changing the type, size, shape, volume fraction, interface, and degree of dispersion or aggregation of the nanofillers enables a great amount of unique combinations of properties with high potential for successful commercial development [2]. In technical literature the terms ‘phase’ and ‘component’ are often used interchangeably. In thermodynamics, however, a clear distinction is made. Thus, a phase is defined as a chemically and physically uniform quantity of matter that can be separated mechanically from a nonhomogeneous mixture. Hence, multiphasic polymer systems would comprise at least two different phases, e.g. semicrystalline single polymers featuring amorphous and crystalline phases or polymorphic polymers containing different concomitant crystalline phases. Many polymer systems with industrial importance such as blends, colloidal polymers, polymer composites/filled polymers, etc. comprise two or more chemically distinct components, each one of them being able to contain various phases as well. These components may have different sizes, with microscopic to nanoscopic blocks being present.

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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The wide use of multicomponent and multiphasic polymer systems in polymeric products fostered the investigations on their structure development during processing and the establishment of structures–properties relationships [3, 4]. Apart from their industrial importance, multiphase polymers and multicomponent materials on their basis are model systems in statistical physics for studying fundamental aspects of many properties such as conformational properties of the chains, the kinetics of phase transitions, as well as the detailed dynamics of diffusion processes [5]. The large molecular dimension of polymer systems markedly reduces the mixing entropy and provides the basis for self-organized structures [6]. Because of all these reasons, investigating polymer systems comprising many components and phases has become an important issue within modern materials science. Generally, scattering methods are a useful tool to study a multiphase or multicomponent system since they are sensitive to the spatial inhomogeneities due to composition or phase fluctuations in polymer materials, either amorphous or semicrystalline. Many relevant studies in this field have been performed by means of small-angle scattering of X-rays (SAXS) or of neutrons (SANS) [7]. The latter technique is less accessible due to the necessity of nuclear reactors and special safety precautions. The wide angle scattering of X-rays (WAXS) called also X-ray diffraction is very frequently used in characterization of semicrystalline multiphase systems. The diffraction pattern contains information that is specific to each phase within the illuminated volume, including both geometric and structural parameters, many of which are inaccessible to other techniques. It is a common feature of all scattering techniques that the structural information can be collected non-invasively, providing in situ and real-time possibilities. While these capabilities already turn WAXS and SAXS into quite powerful techniques, the output can be considerably enhanced by collecting data in synchrotron beamlines, using high-flux and, whenever possible, micro-focused X-ray beams [8]. Although synchrotron is a good option for doing WAXS and especially SAXS, some benchtop equipment is also available (e.g. with a Kratky camera or with a two-dimensional detector as the NanoStar analyzer of Bruker AXS) capable of yielding good SAXS data. Characterization of multicomponent/multiphase polymer systems by X-ray techniques is a vast scientific area. It includes by definition the studies on all polymer blends, alloys and composites that are made of chemically-independent constituents (components). Moreover, each semicrystalline polymer is to be considered at least biphasic, comprising a disordered (liquid) amorphous fraction and a solid fraction with crystalline order. If the latter is made of various crystalline phases (polymorphs) even a semicrystalline homopolymer can be considered a multiphase system. With all these ideas in mind, the scope of the present chapter had to be limited to some recent studies on the application of synchrotron WAXS and SAXS in three particular multicomponent and multiphase polymer systems. The first system comprises materials that became known as microfibrillar reinforced composites (MFC) produced from oriented blends of thermoplastic semicrystalline polymers by conventional processing techniques. These materials belong to fiber-reinforced composites that have many important engineering applications but are notoriously difficult to study [9]. As a second material system, the structure development during processing of an immiscible polymer blend of polypropylene (PP) and polystyrene (PS) was investigated by X-ray scattering techniques. Structure formation in polymers blends has been widely investigated in the last years, mainly in terms of the development of the size, shape, and orientation of the dispersed component under flow deformation [10]. Further, the structure evolution and damage during stretching in the solid state of polymers blends is a much less researched topic. Complementing, this second study, the structure evolution of the PP/PS blend was investigated by time resolved X-ray scattering in a synchrotron source. Finally, the third case reveals investigations on the structure of polymer nanocomposites developed during processing and also during stretching. Polymer nanocomposites are a recent class of materials, and very few studies have been published on the structure development of polymer nanocomposites (e.g. [11, 12]).

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In this chapter, it will be also demonstrated how WAXS and SAXS can be used, along with other characterization techniques, to characterize the structure of multicomponent and multiphase polymer systems and how their nanostructure relates to their mechanical behavior and properties.

17.2 Theoretical Background 17.2.1

Microfibrillar Reinforced Composites (MCF): Definition and Preparation

The MFCs, described initially about two decades ago [13–15], are a special type of in situ nanocomposites combining the easier preparation and processability of the conventional polymer fibrous composites with the presence of micro- and nanosized, high aspect ratio reinforcements typical of the nano- and molecular composites. In the MFCs these reinforcements are fibrils built of bundles of flexible, organic macromolecules produced by appropriate mechanical and thermal treatment of a polymer blend. The typical diameters of the reinforcing fibrils in MFC were found to be within the upper size limit of nanocomposites (i.e. 100–1000 nm). Nevertheless, they could hardly be considered typical representatives of either macro- or nanocomposites [16]. The preparation of MFCs is quite different from that of the conventional composites, insofar as the reinforcing micro- or nanofibrils are created in situ during processing, as is the relaxed, isotropic thermoplastic matrix. The preparation of MFCs comprises three basic steps [17, 18]. First, melt-blending is performed of two or more immiscible polymers with melting temperatures (Tm ) differing by at least 30 ◦ C. In the polymer blend so formed, the minor component should always originate from the higher-melting material and the major one from the lower melting component or could even be amorphous. Second, the polymer blend is drawn at temperatures equal or slightly above the glass transition temperatures (Tg ) of both components leading to their orientation (i.e. fibrillation). Finally, liquefaction of the lower melting component is induced thus causing a nearly complete loss of orientation of the major phase upon its solidification, which, in fact, constitutes the creation of the composite matrix. This stage is called isotropization. It is very important that during isotropization the temperature should be kept below Tm of the higher melting and already fibrillated component. In doing so, the oriented crystalline structure of the latter is preserved, thus forming the reinforcing elements of the MFC. Although MFCs are based on polymer blends, they should not be considered oriented blends. It is the stage of isotropization where the latter are transformed into composite materials. Along with the loss of orientation of the matrix, depending on the chemical functionality of both reinforced and reinforcing components, chemical reactions may also take place resulting in the formation of a copolymeric interface. This interface plays the role of a compatibilizer increasing the adhesion between the matrix and the reinforcing components. If no chemical functionality is present, suppressing of incompatibility between the two materials may be achieved by adding compatibilizing agents to strengthen the interface. In the first studies on MFCs, the composites were prepared on a laboratory scale performing every one of the aforementioned three processing stages separately, one after another. Blending was realized in a laboratory mixer or a single-screw extruder to obtain non-oriented strands that were afterward cold-drawn in a machine for tensile testing, followed by annealing of the oriented strands with fixed ends [13–15, 19–22]. Obviously, this discontinuous scheme is difficult to apply in large-scale production. More relevant in this case are the continuous setups developed more recently [23–26]. Blending of the components and extruding the initial strands could be performed in a twin-screw extruder coupled with two or more drawing devices [24]. After the extrusion blending–drawing stage, one obtains the polymer blend at the exit of the second haul-off device in the form of oriented, continuous cables (OC). To perform the matrix isotropization stage, these strands are further processed by compression molding at temperatures above Tm of the matrix and below

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Tm of the reinforcing fibrils, whereby the former melts assuming the form of the mold and embedding the bundles of oriented fibrils whose orientation and length may be varied [27]. Other molding techniques can also be used instead of compression molding. Monticciolo et al. [28] and later on Pesneau et al. [29] and Evstatiev et al. [30] used an approach in which after the fibrillation of the respective blend by drawing, the oriented strands were cooled down to freeze the morphology and chopped to pellets. The latter were reprocessed by second extrusion or by injection molding at a temperature below the Tm of the dispersed fibrillated component. In this chapter an attempt is made to explain the mechanical properties of polyamide reinforced MFCs with a HDPE matrix relating them to their structure using synchrotron WAXS and SAXS complemented by scanning electron microscopy (SEM). 17.2.2

Clay-containing Polymer Nanocomposites

Polymer nanocomposites exhibit substantial improvement of their properties as compared with those of the virgin raw materials and their traditional microcomposites [11, 12, 31, 32]. Among the latter, nanoclayreinforced polymers received a lot of industrial and scientific interest as they show improved mechanical properties, higher thermal resistance, reduced gas permeability and reduced flammability [11, 32]. The influence of the incorporation of nanoclays on the structure and hierarchical organization of a polymer, and therefore on their properties, has been investigated. When fillers of size comparable with the segmental blocks of macromolecules are added to the polymeric systems they interfere at the molecular level giving unusual properties, at low levels of incorporation (even less than 1%) [31, 33]. For nanoclay-filled polymers, structural features as nanoparticles dispersion and exfoliation level influence the nanocomposite properties [34, 35]. When adding a nanoclay to a polymer three materials may be obtained depending upon the degree of separation of nanoclay layers (Figure 17.1): (i) a fully exfoliated clay, resulting in a true polymer nanocomposite; (ii) intercalated clays with increased intergallery distance due to the insertion of the polymer macromolecules; (c)

0

1

2 2θ (°)

3

4

2θ = 2.72° d1 = 3.25 nm

0

1

2 2θ (°)

3

4

Intensity (a.u.)

d1

d0

(b)

(a)

Intensity (a.u.)

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0

1

2

3

4

2θ (°)

Figure 17.1 Structure of a clay-filled polymer and respective Intensity-2θ plot (WAXS): (a) exfoliated clay nanocomposite; (b) intercalated composite (d1 – clay inter-gallery distance; d1 > d0 ); (c) agglomerated (tactoids) microcomposites (d0 – pristine clay inter-gallery distance).

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and (iii) agglomerated nanoclays, resulting in a traditional polymer microcomposite. In most of the cases, fully exfoliation of the nanoclays is difficult to achieve and the property enhancement is reduced. Due their nanometer size, X-ray diffraction (WAXS) is able to measure the clay’s intergallery distance. However, for very low levels of incorporation of nanoclays, the absence of a reflection peak is not a direct evidence of their fully exfoliation. The deagglomeration/exfoliation and dispersion of the nanometric reinforcements in the polymeric matrix is essential for improved mechanical behavior [32]. Contrasting with the innumerous studies on polymer nanocomposites in the last years, the evolution of the structure during deformation has been investigated scarcely [36]. In this chapter, the structure development during processing of nanoclay-filled PP is investigated. The morphology of PP is characterized for nanocomposites with different amounts of incorporation of nanoclays. Furthermore, the structure evolution during stretching of PET nanocomposites is studied by in situ SAXS investigations. This aims at giving fundamental insights about deformation mechanisms at the nanoscale with adequate time-resolution. 17.2.3

The use of WAXS and SAXS in Characterization of Polymers

˚ in wavelength. Scattering X-rays are electromagnetic radiation occupying the spectrum from 10−2 to 102 A experiments with polymers are performed mostly with the Kα characteristic radiation from a copper target ˚ X-rays of similar wavelength can also be selected, by means of monochromator, tube with λ = 1.5418 A. from the broad spectrum emitted by a synchrotron radiation source. The setup scheme of a scattering experiment in symmetrical-transmission mode frequently used in synchrotrons is represented in Figure 17.2(a). There, radiation from the storage ring source (1) is monocromatized by the incident beam optics (2). The sample (3) is in upright position and the primary (unscattered) beam passes through it hitting the beamstop (5). The detector (4) collects only the X-rays scattered at various 2θ angles. This setup allows the easy change of the sample-to-detector distance R, but requires a two-dimensional detector for recording the complete scattering pattern. Figure 17.2(b) shows a scheme of the classical symmetrical-reflection geometry X-ray setup. With this setup the angle 2θ is changed while recording the intensity of the scattered radiation measured typically by a linear detector. To obtain the complete scattering pattern, the sample should be tilted and/or rotated.

Figure 17.2 (a) – Symmetrical-transmission geometry of the X-ray setup; (b) – Symmetrical-reflection geometry X-ray setup. 1 – X-ray source; 2 – Incident beam optics; 3 – sample; 4 – detector; 5 – beamstop; R – sample-todetector distance.

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Handbook of Multiphase Polymer Systems Table 17.1 Subareas of scattering as a function of the sample-to-detector distance R assuming an X-ray wavelength of λ ≈ 0.15 nm. Reprinted from [8]. Copyright (2007) Springer. Subarea

R, m

Focus

WAXS MAXS SAXS USAXS

0.05–0.2 0.2–1.0 1.0–3.0 6.0–15.0

Arrangements of chain segments Liquid-crystalline structure Nanostructure 3–50 nm Nano and microstructure 15 nm–2000 nm

Scattering experiments in multiphase polymer systems are carried out in four different angular regions (subareas) shown in Table 17.1 [8]. When the symmetrical-transmission geometry is used, switching between the various subareas is made by an arbitrary change of the distance R. The scattering patterns obtained in the WAXS subarea yield information on the arrangement of polymerchain segments e.g., orientation of the crystalline and amorphous phases, crystalline structure, size of crystals, crystal distortions, WAXS crystallinity. The subarea of middle-angle X-ray scattering (MAXS) covers the characteristic scattering of liquid-crystalline structure and rigid-rod polymers. In the SAXS regime the typical nanostructures are observed. Because of the long distance between sample and detector time-resolved measurements can only be carried out at synchrotron radiation sources. The ultra small-angle X-ray scattering (USAXS) extends the accessible structure towards the micrometer range. Time-resolved measurements require a synchrotron beam that is intensified by an insertion device [8]. The most frequently used subareas in scattering experiments with polymers are WAXS and SAXS. Generally, the interactions of X-rays with matter produces two different phenomena [7]: (1) scattering of X-rays by the individual electrons in the sample, and (2) interference among the various scattered waves. Clearly then, the term scattering refers only to phenomenon 1, while the term diffraction is related to the combination of 1 and 2. In fact, this distinction is often omitted. Thus, when the scattering pattern is diffuse and especially if it is in the SAXS subarea, the term scattering is exclusively used, even if some interference is involved. The term diffraction tends to be used when the sample is crystalline, i.e. sufficiently regular so as to concentrate the scattered beam around a number of sharply defined scattering directions, as is in the WAXS subarea of semicrystalline polymers. In principle, the X-ray experiment measures the X-ray flux (i.e. the intensity of the scattered radiation, Is ), as a function of the scattering direction, determined by the scattering vector s or scattering angle, 2θ (Figure 17.3). These data are then analyzed and interpreted so as to obtain information about the relative placements of electrons in the sample. The theory of the X-ray scattering in polymers has been rigorously developed and extensively described in the specialized literature [7, 8, 37, 38]. This chapter will only focus on some specific applications of WAXS and SAXS useful for characterization of multiphase and multicomponent systems with the emphasis on the applied aspects and not on the theoretical and mathematical ones.

17.2.3.1

Degree of Crystallinity by WAXS

The scattering from an amorphous material such as a melt or a glass gives an intensity pattern which is broad and essentially featureless except for the so-called amorphous halo. On the other hand, the diffraction pattern obtained from a crystalline material consists of series of sharp Bragg peaks, easily distinguishable from the diffuse background (Figure 17.2). Consequently, a semicrystalline polymeric material will give

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(110)

Isotropic PP

radial

I (a.u.)

s

(040) (111+041) (130) (117)

integration

|s| = 4π sin(θ)/λ

10

12

14

16

18

20

22

24

26

28

30

2θ (deg)

Oriented PP φ (110)

azimuthal

meridian

equator

integration

I (a.u.)

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0.5

1

1.5

φ (rad)

Figure 17.3 2D WAXS patterns of an isotropic and highly oriented (stretched) PP samples and correspondent radial and azimuthal integration profiles (orientation direction is vertical).

a pattern representing the superposition of both of these features, whereby their relative contributions will reflect the relative amounts of the noncrystalline and crystalline phases present. That is why, to turn diffraction measurements into a quantitative tool for evaluating the degree of crystallinity, it is necessary to separate the observed intensity into crystalline and noncrystalline components. After this is done, the degree of crystallinity Xc can be defined as [6]: ∞ 4π 0 s 2 Icr (s)ds Q cr ∞ = Xc = Q 4π 0 s 2 I (s)ds

(17.1)

where s is the scattering vector, I(s) is the total scattered intensity, Icr (s) is its crystalline component and Q and Qcr are the corresponding scattering invariants. There are several limitations for the practical application of the above relation. First, it will hold strictly only for non-oriented (isotropic) samples. For anisotropic materials the wide-angle scattering must be recorded as a function of both scattering angle and sample orientation in a texture goniometer [8], before the data can be isotropized and the exact values for Q and Qcr can be calculated. Second, experimentally, the intensities I(s) and Icr (s) are available only up to a finite upper limit of s and not to infinity, i.e., there will always be a systematic error in Xc due to the truncation of both integrals. Moreover, the lattice imperfections in the crystalline phase cause that part of the scattering intensity of the Bragg peaks to be diverted to the amorphous halo leading to underestimation of Xc . The problems caused by the data truncation and lattice imperfections are resolved by the method for Xc determination proposed by Ruland [39].

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Although they render theoretically satisfying results, evaluations of crystallinity based on Eq. (17.1) are quite tedious and labor-consuming. Over the years, less accurate but simpler methods have been developed. Whenever series of samples of the same polymer system are being studied, there might be interest in obtaining not the absolute value of Xc but rather of some relative measure of its changes, called also ‘crystallinity index’, xc . A simple phenomenological method for computing of xc is based on Eq. (17.2) [8]: χc =

Acryst. Aam. + Acryst.

(17.2)

where Aam. is the area of the amorphous halo and Acryst. – the summed areas of all Bragg peaks after subtraction of the machine background. Thus, the WAXS curves can be fitted by means of a peak-fitting program using narrow peaks for the Bragg reflections and one or more wide peaks to fit the amorphous halo. If the sample contains more than one crystallographically distinguishable phases, their type and percentage can be studied by analyzing the number and the positions of the respective crystalline peaks that enter in the Acryst . This approach was adopted in the present chapter. A number of approximate methods for crystallinity measurements in polymer systems are reviewed in [38], Chapter 5. 17.2.3.2

Crystal Size by WAXS

In the equatorial scan, the width (2θ ) of a given (hkl) peak is inversely proportional to the average size of the ordered domains. The average size of the ordered domain contributing to that diffraction peak along the direction normal to the (hkl) plane, Dhkl , is given by the Scherrer equation: Dhkl =

0.94λ FWHM hkl cos(θhkl )

(17.3)

where FWHMhkl is the full width at half maximum of (hkl) peak, λ is the X-ray wavelength and θ hkl is angular position of the (hkl) plane (half of the scattering angle, 2θ ). The determination of FHMHhkl must be performed in peaks that do not overlap or after the peak’s deconvolution. However, this latter option may induce relatively high errors in the estimation of crystal size. Moreover, the peak width can be broadened by a distribution of Dhkl on the sample or straining of the crystal lattice, leading to a sub-estimation of the crystal size. 17.2.3.3

Phase Orientation Studies by WAXS

The orientation of the (hkl) reflection planes can be assessed by WAXS. In an oriented semicrystalline polymer, the intensity of each reflection is very dependent upon the azimuthal angle (φ). Assuming a given unit crystal geometry, the intra-chain orientation (00l) is revealed on the meridian and the inter-chain (hk0) on the equator. The more concentrated is the intensity in a reflection arc the higher is the level of crystalline phase orientation. This level of orientation is calculated from azimuthal profiles (between azimuthal angle φ = 0 ÷ π /2 rad or φ = 0 ÷ π ; Figure 17.3). The crystalline phase orientation is normally assessed by the average square of the cosine of φ (the angle between the crystallographic plane and the reference direction), cos2 ϕ, defined as (assuming an uniaxial symmetry):  2  cos φ =

 π/2 0

I (φ) cos2 φ sin φdφ

 π/2 0

I (φ) sin φdφ

(17.4)

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where, I(φ) is the scattered intensity of (hkl) plane as function of φ. The crystalline phase orientation is normally expressed by the Hermans’ orientation function, f , (or second moment of orientation function) defined by:   3 cos2 φ − 1 f = 2

(17.5)

According to Eq. (17.5), if all chains are oriented on the director direction f = 1; for fully transverse orientation f = – 0.5; and f = 0 for random orientation. If the plane normal coincides with a crystallographic axis (a, b and c) and for orthorhombic unit cell geometry, the orientation of the different unit crystal planes ((100), (010) and (001)) is interrelated by:  2      cos φ100 + cos2 φ010 + cos2 φ001 = 1

(17.6)

More generally, if the crystallographic planes are not parallel and other planes of orientation are known, the other f can be calculated by using the Wilchinsky’s method [40]. 17.2.3.4

Structure Investigations by SAXS

As seen from Table 17.1, SAXS is used to study structures of size on the order of 3 nm and larger. Information on such relatively large structures can be collected at scattering angles 2θ lower than 2◦ . The relation between the dimensions in the real space and those in the reciprocal space is given by the Bragg law: d=

λ = s −1 2 sin θ

(17.7)

Hence, if the distance d representing the period of repetition in the structure (e.g. the distance between the similar domains), is around 10 nm and larger, the corresponding scattering angle 2θ will be about 0.6◦ or smaller, i.e. the observations will be made in the SAXS subarea. On the other hand, the typical distances ˚ between the crystallographic planes in polymers are in the order of few Angstroms, therefore the scattering angle will be typically about 20◦ , i.e. in the WAXS subarea. It should be noted that all methods developed for the analysis of WAXS data are applicable in SAXS analysis as well. In the latter case there exist theoretical results dedicated specifically for SAXS data treatment. For example, in small angle scattering sin θ can always be approximated by θ . Moreover, in the SAXS analysis it is assumed that any details of size scale less than 0.1 nm do not exist. In the simplest and most frequently-used analysis of SAXS data, the observed peak of the scattering curve is related to the average distance between the nanoscopic domains in the sample, called also the long period, L. Hence, for small scattering angles and based on the reciprocity in the Bragg law: L = 1/smax

(17.8)

In Eq. (17.8) L represents the sum of the average thickness of the crystal lamellae, lc and of the interlamellar amorphous regions, la . In isotropic and moderately oriented polymer samples smax must be measured only after background subtraction and Lorentz correction of the curve [8], i.e. after changing the y-axis from I(s) to s2 .I(s). Only in this case the L-data obtained directly from the scattering curve maxima are close to the correct ones calculated by more rigorous methods [41]. Highly anisotropic materials that show intensive peaks on a relatively low

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background do not require Lorentz correction. In this chapter, the correct smax values of various multiphase polymer systems were obtained by the above procedure and discussed as a function of the materials processing parameters. Apparently, Eq. (17.8) cannot be used for determination of lc and la . To do that, the approach of Kortleve and Vonk [42] elaborated for the case of isotropic polymers is to be employed. The Fourier transform of the Lorentz corrected SAXS profile is calculated, namely the linear correlation function γ1,r (CF) as: γ1,r = Q

∞ 0

(I − Ib )q 2 cos(qr ) exp(σ 2 q 2 )dq ∞ 2 0 (I − Ib )q dq

(17.9)

Here, q = 2π.s;Ib is that contribution to the total scattering arising from density fluctuations (liquid scattering), and σ is a term, related to the thickness of the crystal/amorphous interface. Q is the so-called scattering invariant that can be determined by integrating the SAXS profile over all scattering angles, i.e.: 

∞

Q=

(I − Ib )q 2 dq

(17.10)

0

The advantage of the CF method in contrast to Bragg’s law is that, in addition to the long period L the values for lc and la and the degree of crystallinity within the lamellar stacks (xcl ) (also called linear crystallinity) can be obtained. In addition to the L value, the CF approach calculates for each sample two additional estimates for the long spacing – from the position of the first maximum of CF (denoted as L cM ) and from twice the position of the first minimum of CF L m c . To calculate the values of la and l c on the basis of CF, the following equation was used: B = x1 (1 − x1 ) L cM

(17.11)

where B is the position of the first intercept of CF with the r-axis. From the two solutions x1,2 of the above quadratic equation, the one with the higher value is ascribed to the larger fraction of the two phases found within the lamellar stacks. For example, in highly crystalline samples, x2 would correspond to the crystal fraction within the lamellar stacks (denoted as xcl ) and 1 – xcl would, then, represent the amorphous fraction within the stack. Once the assignment of xcl is made for each particular case, one may calculate the lc and la from the values of L employing the following equations: lc = xcl L

and la = (1 − xcl )L

(17.12)

M where L, as indicated in [5], may take the values of L m c , or L c . The CF analysis [42] and the more recently developed linear interface distribution function (IDF) [43] are not applicable in anisotropic polymer systems such as MFCs, but work well in clay-filled polymers. Some recent examples of the application of the CF approach in PA6/montmorillonite nanocomposites will be given in Section 17.3.1. Other applications of CF and IDF approaches in isotropic multiphase polymers can be found in [44–46].

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17.3 Studies on Multiphase Polymer Systems 17.3.1

Polyamide 6/montmorillonite Nanocomposites

It is important to note that in these two-component systems the SAXS patterns can be treated quantitatively to extract two types of structural information: related to the polymer matrix and related to the inorganic filler. Here, the study of the polymeric part of the composite will be exemplified. Some recent studies on the SAXS patterns due to clay are also available [47–49]. 17.3.1.1

Materials and Sample Preparation

This series was prepared by extrusion-blending of neat PA6 with a PA6/nanoclay masterbatch containing 20% of montmorillonite (MMT) supplied by Nanocor. After compression molding of the respective granulates, isotropic PA6 nanocomposites in the form of laminate plates were obtained in which the MMT content varied between 1 and 7.5 wt.%. 17.3.1.2

Experimental Techniques

The structure of the PA6/MMT laminates was studied by SAXS in the setup previously described. The 2D SAXS patterns were cut to obtain the respective 1D scattering profiles which were then analyzed by the SASDAP software (Copyright © 1995 by R. Verma, A. Biswas and B. Hsiao, DuPont Experimental Station, M Wilmington, DE, USA). This software computes the linear CF and derives values for L m c , L c , la and lc according to Eqs (17.8–17.12). 17.3.1.3

Structure of PA6/MMT Composites

Figure 17.4 shows the dependence of the Bragg long spacing LB on the MMT content. It can be seen that with the increase of the MMT content, the intensity of the PA6 peak decreases until it completely disappears

Intensity I, a.u.

1

5 2 3

4

0,02 0,04 0,06 0,08 0,10 0,12 0,14 0,16 0,18 0,20 Scattering vectors, nm–1

Figure 17.4 1D SAXS profiles of PA6/MMT nanocomposites containing various amounts of montmorillonite clay, wt.%: 1 – no MMT, 2 – 1%; 3 – 2.5%; 4 – 5%; 5 – 20% (Nanocor masterbatch).

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Handbook of Multiphase Polymer Systems Table 17.2 Comparison of structural parameters obtained by WAXS and SAXS as a function of the MMT content in PA6/MMT nanocomposites. For designations see the text. Sample composition + parameters WAXS

Structural parameters (SAXS)

PA6 no MMT Crystallinity index: 0.41 γ = 0.168; α = 0.242; α/γ = 1.44

LB = 90 A˚ ˚ LcM = 91 A

˚ l1 = 58 A ˚ l2 = 33 A xc = 0.664

PA6 1% MMT Crystallinity index: 0.43 γ = 0.290; α = 0.140; α/γ = 0.48

LB = 93 A˚ ˚ LcM = 93 A

˚ l1 = 59 A ˚ l2 = 34 A xc = 0.616

PA6 2.5% MMT Crystallinity index: 0.335 γ = 0.231; α = 0.194; α/γ = 0.84

LB = 104 A˚ ˚ LcM = 97 A

˚ l1 = 63 A ˚ l2 = 34 A xc = 0.656

PA6 5.0 % MMT Crystallinity index: 0.388 γ = 0. 268; α = 0.120; α/γ = 0.45

LB = 120 A˚ ˚ LcM = 102 A

˚ l1 = 68 A ˚ l2 = 34 A xc = 0.629

PA6 20% MMT Crystallinity index: 0.335 γ = 0.261; α = 0.074; α/γ = 0.28

LB = 132 A˚ ˚ LcM = 117 A

˚ l1 = 83 A ˚ l2 = 34 A xc = 0.667

in the PA6/5% MMT sample. It seems that the introduction of MMT decreases the density difference between the amorphous and crystalline PA6 phases. In the sample with the largest amount of MMT, it is the amorphous phase that becomes denser. The reappearance of the density difference results in a new SAXS peak corresponding to larger LB (Table 17.2). As a result of this observation one may conclude that, if a simple two-phase model is considered for the PA6 matrix, the nanoclay is distributed within the amorphous phase and not within the crystalline one. Figure 17.5 shows the CF curves of the PA6/MMT composites compared to that of the neat PA6 matrix. Curves 4 and 5 reveal larger long spacing values (denoted in Table 17.2 with LcM to distinguish from those obtained from the raw SAXS profiles), with broader size distributions, as compared to the CFs of samples 1–3. The CF analysis allows the division of LcM into two parts: l1 (larger) and l2 (smaller) values, corresponding to either the average thickness of the crystal lamellae, lc , or to that of the interlamellar amorphous layers, la . Table 17.2 contains also the values of the intra-stack crystallinity fraction denoted as xc . As explained in the notes to Eq. (17.11) above, the value of xc is supposed to be always higher than the average crystallinity index, obtained by WAXS. The assignment of the l1 and l2 values is only possible after consideration of all data from WAXS and SAXS in the table. Thus, with the increase of the MMT content there is a general trend toward a slight diminution of the average WAXS crystallinity index, whereby the fraction of the γ -PA6 polymorph notably increases. At the same time, both LB and LcM values increase, due to the augmentation of the larger size l1 , whereas the l2 values remain constant. Having in mind the explanation of the intensity changes of the SAXS peak in Figure 17.4, it can be concluded that the modification of PA6 with MMT results in the expansion of the amorphous phase, in which the MMT is concentrated, i.e. l1 = la and l2 = lc . Forthcoming studies in this PA6/MMT system will allow correlating the said nanostructural changes with the mechanical properties of the respective nanocomposites.

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–0.4 γ1,r /Q

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150 200 r, in Angstroems

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Figure 17.5 Linear correlation function curves of PA6/MMT nanocomposites containing various amounts of montmorillonite clay, wt.%: 1 – no MMT, 2 – 1%; 3 – 2.5%; 4 – 5%; 5 – 20% (Nanocor masterbatch).

17.3.2 17.3.2.1

Microfibrillar Composites (MFC) Materials and Sample Preparation

For the preparation of the oriented MFC precursors, the selected amounts of the HDPE, PA6 and YP granulates were premixed in a tumbler in various proportions. Each mixture was introduced into a K-Tron Soder gravimetric feeder that fed it to the hopper of a laboratory modular Leistritz LSM 30.34 intermeshing co-rotating twin-screw extruder [50]. At the exit of the last haul-off device the blends are in the form of oriented, continuous cables. The OCs were cut and aligned in the form of unidirectionally arranged bundles (unidirectional ply laminate, UDP. This was subjected to selective melting whereby isotropization and controlled crystallization of the matrix occurred in a hot press at a fixed temperature of 160 ◦ C, a pressure of 2 MPa and a cooling rate of ca. 10 ◦ C/min. Standard rectangular laminate plates (60 × 120 mm with a thickness of 0.1–1.5 mm) were obtained from all the precursors. They were used for structural and morphological characterization, as well as to yield specimens for the tensile tests.

17.3.2.2

Experimental Techniques

Scanning Electron Microscopy (SEM) To analyze the morphology of the MFCs and their precursors, SEM of freeze-fractured specimens was used. For each blend, specimens were collected for morphological analysis typically at three different locations of the extruder line: at the extrusion die, after the first and after the second haul-off units. The final MFCs obtained after compression molding were also analyzed. All samples were sputter-coated with gold and observed in a Leica S360 SEM at magnifications of × 2.0 k, × 5.0 k and × 7.5 k. Most of the specimens studied were obtained by cryogenic fracture with liquid nitrogen and the fractured surfaces were observed by SEM.

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X-ray Measurements Details All WAXS and SAXS patterns in this study were registered at the Soft Condensed Matter Beamline (A2) of HASYLAB, Hamburg, Germany using synchrotron radiation with a wavelength fixed to 0.15 nm, using the symmetrical transmission geometry (Figure 17.2(a)). The sample-todetector distance for SAXS was set at 2830 mm, the diffraction patterns being registered by means of a MAR CCD 2D detector with exposure times of 30 s. For the WAXS measurements the detector was positioned at 90 mm in respect to the sample. The various MFCs were studied in transmission mode, the sample thickness being in the 0.1–1.5 mm range. Scattering patterns were obtained at certain temperatures in the 30–300 ◦ C range employing a typical heating rate of 20 ◦ C/min. A specially designed sample holder was used allowing for a controlled heating/cooling of the sample in the 25–300 ◦ C range. An IMAGO multi-channel process and program controller of JUMO GmbH & Co. KG was used to regulate the sample temperature in heating or cooling at various rates. The difference between the read-out and real temperature of the sample was found to be 3–4 ◦ C at a heating or cooling rate of 20 ◦ C/min.

17.3.2.3

Results and Discussion

SEM Microscopy of HDPE/PA6/YP MFC The SEM images of the final MFCs (Figure 17.6, 1–4b) confirm that: (i) the PA6 reinforcing component has well-expressed fibrillar morphology, and (ii) the average diameter of these fibrils is in the upper nanometer – lower micrometer range, e.g. from 0.6 to 1.5 μm (samples without compatibilizer) and from 0.5 to 1.0 μm (compatibilized samples). Therefore, depending on the compatibilizer 1aAfter extruder

1b

5 μm

MFC UDP 5 μm

80/20/0 1a

1b

2a

2b

3a

3b

4a

4b

77.5/20/2.5

75/20/5.0

70/20/10

Figure 17.6 SEM images of various HDPE/PA6/YP materials after cryogenic fracture at the various stages of the MFCs preparation (compositions given in wt.%): non-oriented blend right after the extruder die (1–4 (a)); UDP composites fractured along the fibrils axis (1–4 (b)).

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Figure 17.7 2D SAXS images of two HDPE/PA6/YP UDP MFC with compositions: 1 – 80/20/0; 2 –70/20/10; at different temperatures: (a)– pattern of starting MFC at 30 ◦ C; (b) – pattern at 160 ◦ C, heated in the beam; (c) – pattern at 30 ◦ C after heating at 160 ◦ C. The fibril axis is horizontal [47].

content, the resulting composites can be considered as either nanostructured (NPC) or microfibrillar (MFC). Further on the second abbreviation the second abbreviation will be accepted. None of the images of MFCs in Figure 17.6 permits measuring directly the fibrils’ lengths. Their average lengths and aspect ratios of the reinforcing PA6 fibrils could be evaluated indirectly as indicated in [35]. First, the diameters of the PA6 globules embedded in the HDPE matrix are measured in the SEM micrographs of HDPE/PA6 blends before orientation (Figure 17.6, 1–4(a)). Then, the diameters of the fibrils in the final MFCs are measured from the respective SEM images and averaged. Supposing that the fibrils’ geometry is cylindrical and that they are produced by deforming the respective PA6 spheres without the formation of voids, i.e., that the volume of the PA6 nodules at the extruder die (Figure 17.6, images 1–4(a)) and of the respective MFC fibril (images 1–4 (b)) are the same, the average length and the aspect ratio of the latter are computed. Thus, in non-compatibilized PA6-reinforced MFCs the length of the reinforcing fibrils can reach 120 μm, whereas the maximum length in the presence of compatibilizer is ca. 40 μm. This would give aspect ratios of 80–200 and 40–80, respectively. Discussing the MFCs in Figure 17.6 it is worth noting that the fibril thickness in images 1–4 (b) varies as a function of the sample composition. The question arises if the fibrils observed are of pure PA6 or also include at their interface physically or chemically bonded oriented HDPE. This question can be elucidated by X-ray scattering experiments. 2D SAXS Studies of HDPE/PA6/YP MFC It can be supposed that after the matrix isotropization stage the final composite will contain fibrillar reinforcement components embedded in a fully isotropic matrix. The presence of fibril-shaped phase is undoubtedly proved by the SEM micrographs of the UDP MFCs in Figure 17.6. Figure 17.7 represents the SAXS patterns of two HDPE/PA6/YP UDP MFC compositions: without compatibilizer (1) – 80/20/0 and with compatibilizer (2) – 70/20/10 at different temperatures. The first examination of the 2D SAXS patterns (images 1(a) and 2(a) show that both composites contain isotropic scattering of randomly distributed lamellar structures and equatorial scattering maxima attributable to lamellar crystals oriented parallel to the horizontal fiber direction. The isotropic ring and the oriented maxima display

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Figure 17.8 Azimuthal distribution of the scattered intensity in the 2D SAXS images of two HDPE/PA6/YP UDP MFCs: (a) 80/20/0; (b) 70/20/10. 1 – Initial MFC at 30 ◦ C; 2 – in beam heating at 160 ◦ C; 3 – at 30 ◦ C after heating to 160 ◦ C The dashed line indicates the fiber direction [47].

˚ This is a clear indication that the observed oriented reflections cannot similar long spacings of >220 A. ˚ [46]. Consequently, originate from the reinforcing PA6 whose L B values are typically between 70 and 90 A it can be supposed that a fraction of the HDPE matrix material has crystallized upon the oriented PA6 fibrils thus forming a transcrystalline layer (TCL) at the interface. Without a special treatment it is impossible to observe at the same time the HDPE and PA6 scattering in patterns 1(a) and 2(a) because of the strong differences in the scattering intensities. Heating the two UDP MFC samples at 160 ◦ C eliminates the HDPE scattering and reveals the oriented PA6 reflections (images 1(b) and 2(b)). Cooling back to 30 ◦ C causes the HDPE matrix to recrystallize. This process takes place differently in the two MFCs under investigation. The oriented HDPE TCL in 70/20/10 MFC at 30 ◦ C after the selective melting of the matrix maintains its equatorial orientation (Figure 17.7, 2(c)), whereas in the 80/20/0 system it rotates by 90◦ and appears at the meridian (Figure 17.7, 1(c)). Isotropic scattering was also present in the two patterns. This reorientation of the HDPE scattering of is better observed if azimuthal cuts of the above patterns are performed (Figure 17.8). The curve of the noncompatibilized sample (Figure 17.8(a)) clearly shows that after recrystallization the peak of intensity is not at 0◦ (i.e. along the fiber axis) but at –90 or 90◦ . In the compatibilized sample (b) the azimuthal distribution of scattered intensity remains the same at 30 ◦ C and at 30 after 160 ◦ C. It is noteworthy that this reorientation of the lamella that takes place in the noncompatibilized samples is not accompanied by chain direction reorientation, i.e. the chain direction of PA6 and that of the oriented HDPE fraction continue to coincide, as in the starting image at 30 ◦ C. This effect will be discussed in the next section dedicated to the WAXS studies. To make a distinction between the two fractions of HDPE, the subtraction procedure described by Nogales et al. [51] was used. The 2D WAXS patterns were first corrected for the incident beam intensity and then the empty chamber scattering was subtracted. It was assumed that the total scattered intensity could be separated into two contributions: (i) the isotropic contribution from the amorphous chains and the non-oriented crystals, being directly proportional to the azimuthally independent component of the total scattered intensity, and (ii) the oriented contribution from all oriented (with varying degree of orientation) scatterers calculated

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X-ray Scattering Studies on Multiphasic Polymer Systems (a)

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=

Figure 17.9 Deconvolution procedure of the SAXS pattern of the 75/20/5 UDP MFC. (a) – original SAXS image; (b) intensity pattern of the isotropic scattering; (c) intensity pattern of the oriented scatters obtained by subtraction (a) – (b) [12]. The fiber axis is horizontal [47].

by subtracting the azimuthally independent component from the total scattered intensity. To determine the azimuthally independent intensity and to perform the said subtraction, a subroutine incorporated into the POLAR 2.7.1 X-ray software was used [52]. Thus, Figure 17.9(a) shows the pattern of the total scattering of the 75/20/5 MFC composition at 30 ◦ C. The computer-generated 2D image of the isotropic intensity is presented in Figure 17.8(b), and the resulting image obtained after (a) – (b) subtraction, corresponding to the oriented scatterers is shown in Figure 17.10(c). As seen from the latter, the said procedure not only separates the two HDPE components, but also reveals clearly the oriented PA6 fraction located along the equator. In Figure 17.10, a 3D visualization of the initial pattern (a) and that of the oriented scattering (b) for the same 75/20/5 composition is given. Image (b) shows better the PA6 contribution to the oriented part of the scattering, pointed by the arrows. Table 17.3 contains the HDPE and PA6 L B values determined from the scattering patterns of six UDP MFCs with different HDPE/PA6/YP compositions. It can be seen that in the absence of compatibilizer, there are no significant differences between the long spacings values of HDPE lamellae located in the bulk (isotropic) and those of the oriented HDPE lamellae in the transcrystalline layer (oriented). Introducing YP compatibilizer results in smaller long periods in the oriented HDPE fraction, while that of the bulk matrix fraction remains as in the noncompatibilized compositions. Only in the 65/30/5 UDP MFC the distance between the oriented

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Figure 17.10 3D SAXS patterns of UDP MFCs before (left) and after (fight) the subtraction of the azimuthally independent component of the total scattered intensity. The white arrows indicate the scattering of the PA6 reinforcing phase.

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Table 17.3 Long spacing values of the HDPE/PA6/YP UDP composites at 30 ◦ C without (L B ) and with (L ∗B ) deconvolution. Reprinted from [50]. Copyright (2010) with permission from John Wiley & Sons. ˚ L B, A

L B *, A˚

N◦

HDPE/PA6/YP composition

PA6

HDPE total

HDPE-iso

HDPE-orient.

PA6 orient.

1 2 3 4 5 6

90/10/0 80/20/0 77.5/20/2.5 75/20/5 70/20/10 65/30/5

100.5 90 94 94 87 82

223 229 221 220 215 236

224 225 (231) 224 224 225 (245) 223

222 222 (225) 211 213 210 (214) 231

95 86 91 92 88 77

Note: The values in parentheses were obtained after recrystallization of the HDPE by in beam heating to 160 ◦ C followed by cooling down to 30 ◦ C. [47].

HDPE lamellae is bigger than that of the isotropic fraction. Most probably, this could be explained as a result ˚ of a higher amount of PA6 in this composition. As regards the PA6 L B values, they vary in the 77–95 A ˚ interval. The PA6 long period of 77 A in the 65/30/5 composition is the closest to the value of the neat oriented PA6 [46]. As mentioned above, after recrystallization, the HDPE fraction in the noncompatibilized and compatibilized samples orients in different ways – in the first case the scattering maxima appeared on the meridian, while in the second maintained their position on the equator. As seen from Table 17.3, in both 80/20/0 and 70/20/10 samples, an increase of the L B values of the isotropic HDPE – in the presence and in the absence of compatibilizer – was observed after matrix recrystallization (the data in parentheses). A significant improvement of the amount and quality of structural information extracted from the oriented part of the scattering in the UDP SAXS patterns (Figure 17.8(c)) can be achieved by image reconstruction followed by computation of the respective Chord Distribution Function, CDF, as suggested in a series of recent papers by Stribeck et al. [53, 54]. An example of this treatment is given in Figure 17.11 for the two UDP MFC samples without and with compatibilization. The image reconstruction comprises background correction and calibration for beam intensity, followed by filling of the ‘blind’areas behind the beam stop and its wire. Thereafter, the computation of the CDF is performed. It is worth noting that all these procedures are R programming environment [55]. carried out automatically using the pv Wave Analyzing the reconstructed patterns of the oriented scattering in the two samples under investigation, it can be concluded that the respective LB values of HDPE (oriented and isotropic) as well as of the PA6 almost coincide with those in Table 17.3. Interesting information can be extracted from the quite complex CDF functions. It is important to note that the CDF images reflect mostly the structure of the HDPE-oriented TCL and to a lesser extent that of the PA6 fibrils. Moreover, the CDF reveals more features than the respective SAXS pattern, being its transformation to the real space. Thus, in the as-prepared UDP MFC without compatibilization (80/20/0 sample), the HDPE TCL forms a microfibrillar system with first and second order long periods, with the HDPE domains being side-by-side and not shifted. Introducing compatibilizer (70/20/10 sample) maintains these structural features but results in the appearance of tilted oriented HDPE domains. The differences between the oriented SAXS fractions in compatibilized and noncompatibilized UDP MFC materials and especially in the corresponding oriented precursors become better pronounced during a simultaneous SAXS/straining experiment. As it was demonstrated recently [55], the automated procedures described above allow the processing of hundreds of data frames associating the structural features with the simultaneously obtained mechanical behavior in stretching or load cycling modes.

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X-ray Scattering Studies on Multiphasic Polymer Systems HDPE/PA6/YP=70/20/10

Chord Distribution Functions (CDF)

Reconstructed

HDPE/PA6/YP=80/20/0

687

Figure 17.11 Reconstructed SAXS patterns (oriented scattering) of two UDP MFC materials and their respective chord distribution functions. Fibril axis is vertical. The CDF function is presented in absolute values (both positive and negative faces in one image). For more details see the text.

2D WAXS Studies of HDPE/PA6/YP MFC The SAXS studies of UDP MFC materials gave evidence that the reinforcing fibrils most probably have a layered, coaxial structure: a core of oriented PA6 and a shell of oriented, transcrystalline HDPE. The WAXS measurements supported this hypothesis. The visual inspection of the 2DWAXS patterns of UDP MFCs (Figure 17.12) shows that the crystallographic characteristics of HDPE and PA6 are very similar leading to a strong overlapping of the respective diffraction peaks. Nevertheless, one can notice that at 30 ◦ C there is a co-existence of isotropic Debye rings and crystalline reflections oriented parallel to the horizontal fibril direction. At 160 ◦ C the HDPE reflections change into a diffuse amorphous halo revealing the oriented PA6 reflections. To separate the contribution of the isotropic and oriented crystalline fractions and to study their origin, the same subtraction procedure as with the SAXS patterns was applied. Figure 17.13 exemplifies this treatment for the 80/20/0 (a) and 70/20/10 (b) HDPE/PA6/YP UDP MFCs showing the starting real 2D WAXS patterns (left), the computer-generated isotropic part of the scattered intensity (center) and the resulting 2D WAXS images of the oriented part after subtraction (right). Subtracting the isotropic crystalline and amorphous fractions allows the outlining of the oriented crystalline reflections that are otherwise undetectable. Together with the expected oriented PA6 reflections in the right images in Figure 17.13, one observes also clear reflections of the oriented matrix. The two weak equatorial arcs belong to the (200) and (002/202) planes of PA6 and the other two, more intense equatorial reflections belong to the (110) and (200) planes of orthorhombic unit cell of HDPE. This is one more indication for epitaxial crystallization of matrix material upon the reinforcing fiber, whereby the chain direction in the matrix crystals coincides with that in the reinforcing PA6 fibrils. Judging from Figure 17.13, this observation is valid for both selected samples – noncompatibilized (a) and compatibilized (b).

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Figure 17.12 2D WAXS patterns of HDPE/PA6/YP microfibrilar composites taken at various temperatures. Fibril direction is horizontal [47].

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–

Figure 17.13 Example of the analysis of the WAXS patterns at 30 ◦ C of UDP MFCs: Left – total scattered intensity; Center: calculated isotropic intensity; Right: oriented scattered intensity. (a) – 80/20/0 and (b) – 70/20/10. The fiber axis is vertical [47].

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Figure 17.14 3D WAXS patterns of UDP MFCs before (left) and after (right) the subtraction of the azimuthally independent component of the total scattered intensity. The white arrows point at the (200) reflections of α PA6 [47]. Reprinted from [50]. Copyright (2010) with permission from John Wiley & Sons.

Figure 17.14 shows the 3D images of the real WAXS patterns before treatment (left) and of the oriented scattering after subtracting (right) of the same two MFCs. The white arrows indicate the position of the α-PA6 (200) reflection. This representation shows better the anisotropy of the HDPE (110) and (200) diffractions. For a quantitative evaluation of oriented and isotropic parts of the total scattered intensities, the respective 2D WAXS patterns were integrated in the 0–180◦ range to get the 1D WAXS profiles, which were afterwards fitted by Gaussian peaks. The results from peak-fitting applied in the 80/20/0 MFC sample are presented in Figure 7.15(a) and (b). The deconvolution of the integral profile of the oriented part clearly shows the (110), (200) and (210) contributions of the HDPE (Figure 17.15(a), the shaded reflections)) and also the four crystalline reflections of α- and γ PA6. The peak-fitting of the isotropic part displayed crystalline reflections (110), (200) and (210) of the HDPE matrix only and the amorphous halos of PA6 and HDPE, respectively (Figure 17.15(b)). Based on the angular position of the reflections, the d-spacings (dhkl ) of the corresponding planes were calculated. A quantitative evaluation of the peak-fitting results for two representative MFCs – without (80/20/0) and with compatibilization (70/20/10), as well as data for d-spacings are given in Table 17.4.

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Handbook of Multiphase Polymer Systems Table 17.4 Results from the deconvolution of the oriented and isotropic part of 2D WAXS patterns of selected HDPE/PA6/YP UDP MFC. HDPE/PA6/YP 80/20/0 WAXS Reflections

2θ, deg.

70/20/10

Content, %

˚ dhkl , A

2θ, deg.

Content, %

˚ dhkl , A

28.7 7.6 38.2 7.6 6.9 9.1 1.9

4.34 4.07 4.05 3.99 3.76 3.65 2.95

Oriented part of WAXS intensity (200) – α PA6 (001) – γ PA6 (110) – HDPE (200) – γ PA6 (002)/(202) – α PA6 (200) – HDPE (210) – HDPE

19.90 21.05 21.44 21.79 23.09 23.69 29.61

28.5 6.6 34.9 13.7 6.9 7.9 1.5

PA6 fraction, % HDPE fraction, %

4.34 4.11 4.03 3.97 3.75 3.65 2.94

19.92 21.35 21.33 21.66 22.99 23.74 29.50

55.7 44.3

50.8 49.2

Isotropic part of WAXS intensity (110) – HDPE (200) – HDPE (210) – HDPE

21.13 23.56 29.29

14.6 11.4 1.9

4.09 3.67 2.96

20.97 23.48 29.24

9.8 12.6 1.3

4.12 3.69 2.97

Notes: In the isotropic part of the WAXS intensity the crystalline reflections only are included. The difference to 100% will give the content of the amorphous HDPE and amorphous PA6. dhkl is the dspacing of the respective crystalline plane. The oriented reflections are considered 100% crystalline [47].

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Figure 17.15 1D W\AXS profiles of the 80/20/0 HDPE/PA6/YP UDP MEF exemplifying the peak-fitting of the oriented scattering (a) and of the isotropic EWAXS scattering (b). The pattern in (a) was obtained after subtraction of (b) from the initial WAXS pattern with the total scattered intensity [47]. Reprinted from [50]. Copyright (2010) with permission from John Wiley & Sons.

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2R 2 = 750 nm

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R2 – R1 = TCL

2R1 Figure 17.16

2R1

Schematic presentation of the fiber cross-sections of 80/20/0 and 70/20/10 UDP MFCs.

From Figure 17.15 and Table 17.4 it can be seen that a significant part of the HDPE matrix is able to crystallize oriented along the PA6 fiber, thus forming a transcrystalline layer in such a way that the chain directions of the two polymers coincide. The rest of the matrix, situated in the bulk, crystallizes isotropically. The relation between the content of the PA6 fibrils and the oriented part of the HDPE matrix (the crystalline fraction) is almost 1.03:1.00 in the 70/20/10 MFC and 1.26:1.00 in the 80/20/0 system. This means that in the presence of compatibilizer a larger part of the HDPE is included in the transcrystalline layer without changing considerably its crystallographic characteristics. Based on the d-spacing values it can be concluded that the HDPE unit cell is slightly larger in the bulk, as compared to that in the oriented transcrystalline layer. The data about the PA6 and HDPE fraction in the oriented scattering can be used to obtain an estimate of the TCL thickness in UDP MFC materials. From the SEM studies (Figure 17.6) one can estimate the average visible thickness of the reinforcing fibrils in the MFC composition. Let us take samples 1(b) and 4(b) and suppose that the fibrils are cylindrical with a PA6 core (with a diameter 2R1 ) being uniformly coated by coaxial transcrystalline layer of HDPE whose thickness is given by R2 – R1 . Therefore, the visible diameter of the fibril estimated form SEM will be R2 . Figure 17.16 gives a schematic view of the cross-sections in the two selected UDP MFCs – without and with compatibilization. Whenever X-rays are interacting with matter, their main partners are the electrons in the studied sample. Thus X-ray scattering is probing the distribution of electron density, ρ(r), inside the material. In the field of WAXS, ρ(r) is identical to the average electron density, ρ. For a given material or specific phase within a material, ρ is calculated as [8]: ρ = NA

ZM ρm MM

(17.13)

Here, ρ m , is the respective average mass density, NA is the Avogadro’s number (6.022 × 1023 mol−1 ), ZM is the number of electrons per molecule or monomer unit and MM – the molecular weight of molecule or monomer unit. Logically, the intensity of the radiation diffracted by either the PA6 or HDPE component will be proportional to the volume of this phase Vi and the respective average electron density ρ i : I Si ∼ Vi · ρi

(17.14)

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If we denote by V PA6 the volume of the PA6 core, it can be written that VPA6 = π L R12 and   VTCL = π L R22 − R12

(17.15) (17.16)

Combining Eq. (17.14) with (17.15) and (17.16), the following simple dependence can be deduced between the visible by SEM fibril radius R2 and that of the PA6 core R1 :  R12

=

R22 .

f k+ f

(17.17)

I PA6

PA6 s and f = I HDPE . wherein k = ρρHDPE s Table 17.5 summarizes the structural information related to the reinforcing fibrils as revealed by SEM and WAXS methods (i.e. 2R1 , 2R2 and R2 – R1 ) for MFC materials without and with compatibilization reinforced by PA6 or PA12. The same table contains also some mechanical data obtained with these composites, as well as with the neat HDPE matrix material and with the neat oriented polyamides. It can be concluded that the formation of transcrystalline layers TCL is a common feature for all MFCs containing either PA6 or PA12. There can be a significant difference between the TCL thicknesses in PA6 and PA12 reinforced composites, as well as in the compatibilized and noncompatibilized MFCs with the same reinforcement. Compatibilization results in thinner fibrils in which not only the polyamide core, but also the TCL are finer. In PA6 reinforced MFCs the TCL is more than two times thicker than in similar HDPE/PA12/YP composites. Obviously, the TCL thickness is directly related to the mechanical performance of the MFCs, whereby the larger the thickness, the lower the properties. Thus, no matter that the E1 value of neat oriented PA6 is much higher than that of PA12, in both compatibilized and noncompatibilized MFCs this values is either similar or lower in the HDPE/PA6/YP materials. At the same time, the σ y of the HDPE/PA12/YP materials are significantly higher, irrespective of the almost coinciding values of the neat oriented polyamides. It is to be noted the superior flexural stiffness of the PA12 composites, which may have to do with the increased flexibility of PA12 and its better compatibility with HDPE. One has to bear in mind also that in the 80/20/0 samples the formation of TCL can be attributed to physical interactions at the matrix–fibril interface. In the 70/20/10 systems, however, it should be a result of chemical reactions between the maleic anhydride of YP

Table 17.5 Dependence between the structural parameters related to the oriented reinforcing component and the mechanical behavior in various MFC materials. HDPE/PA6/YP

2R2 , nm 2R1 , nm R2 -R1 , nm E1 , MPa σ y , MPa CR , MPa

HDPE/PA12/YP

Neat materials

80/20/0

70/20/10

80/20/0

70/20/10

PA6 orient.

PA12 orient.

HDPE isotropic

1000 733 134 1095 57 2624

750 524 113 920 37 2294

700 567 66 1054 64 3414

560 457 52 972 55 3404

– – – 3180 230 –

– – – 2240 233 –

– – – 827 26 1478

Notes: E1 is the secant modulus determined at 1% strain; σ y , is the maximum stress at break and CR is the three point support flexural stiffness determined according to Nunes et al [56].

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and the amide groups of the polyamide [57]. It can be expected that in the latter case the TCL will include polyolefin component from the YP compatibilizer, which is different from the bulk matrix HDPE. This could be one of the possible explanations of the inferior mechanical properties of the compatibilized samples. 17.3.3 17.3.3.1

Immiscible Polymer Blends Materials and Sample Preparation

A blend of polypropylene/polystyrene (PP/PS) was selected for the investigations, comprising a semicrystalline and amorphpous polymers. The two polymers PP/PS were mechanically blended in a tumbler mixer, in the ratio of 70/30 wt%. The blend was then directed injection molded in the form of small tensile specimens. High back pressure was used to promote a better mixing between both components. 17.3.3.2

Experimental Techniques

The structure of the injection molded PP/PS specimens was characterized by SEM and WAXS techniques. The specimens were also stretched in a tensile machine until different strain levels (of 5, 10, 15, 20 and 25%) in the plastic regime (inducing permanent deformation, but allowing molecular relaxation). The structure of the deformed specimens were then analyzed by 2D WAXS and SAXS experiments (not simultaneously), which allowed the characterization of .the structure evolution of the semicrystalline polymer (PP) and damage occurring by nanocavitation phenomena (e.g. at the polymer interfaces) upon stretching. The sample-to-WAXS detector distance was of 13.4 cm and the sample-to-SAXS distance of 280 cm. 17.3.3.3

Results and Discussion

Structure Development During Processing Figure 17.17 shows the solid-state morphology of the PP/PS injection molded blend. In Figure 17.17(a) are presented the polarized light microscopy (PLM) and scanning electron microscopy (SEM) images of the blend. The injection molded specimen evidences a typical skin–core microstructure. In the skin layer, the disperse component (PS) features very thin and highly elongated strands (in the flow direction, FD) due to the high deformation rates applied during molding. In the core, this disperse component is constituted by particles of different shapes (e.g. spherical, elongated fibers) of higher dimensions, mostly oriented transversely to FD [58]. Figure 17.17(b) shows the WAXS and SAXS patterns of the blend. These are average patterns through halfthickness of the molded specimens. From the WAXS pattern, the molded PP/PS blend features an oriented PP crystalline phase, expectantly in the skin region. A bimodal orientation distribution of the crystalline phase of PP is revealed, with a relatively high a*-axis oriented component [59]. This may arise from the development of transcrystalline structures growing on the highly oriented PS component. The SAXS pattern reveals a typical shish-kebab structure, with lamellae perpendicular to FD and growing in lamellar stacks along FD. The central black region of the SAXS patterns suggest that the PP/PS blend may show some nanosized voids, which may be originated at the interface between the two polymers. Structure Evolution During Stretching The structure evolution and damage during solid-state stretching of immiscible (noncompatibilized) PP/PS blend was also investigated by X-ray scattering techniques. Compared with the neat PP, the PP/PS blend shows substantially reduced deformation capabilities. Figure 17.18 presents the obtained stress–strain curve, and correspondent WAXS and SAXS patterns at given strain levels.

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SEM flow direction PLM skin

skin

core

(a)

core

(b)

Figure 17.17 Morphology of the injection molded PP/PS blend: (a) polarized (left) light and scanning electron (right) microscopy images; (b) WAXS (left) and SAXS (right) patterns.

Figure 17.18

Stress–strain curve of PP/PS blend and WAXS and SAXS patterns at given strain levels.

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The PP matrix shows little evolution of its crystalline structure during stretching, with a slight increase on the level of crystalline phase orientation in the stretching direction, SD. This appears to take place at low strains during the initial linear zone of the stress–strain curve (up to ε = 0.05). The crystalline lamellar structure is progressively and partially destroyed on the course of deformation (reduction on the intensity of the two vertical lobules of the SAXS patterns). Significant voiding also occurs during deformation. Voids seem to be nucleated parallel to SD (ε = 0.05), presumably at the interfaces between the two components. Then they increase laterally, growing mainly perpendicular to SD (ε = 0.10). At larger strain levels these voids become bigger and more elongated perpendicular to SD (ε = 0.15). In the meanwhile, the void size increases so that SAXS could not detect them and the SAXS pattern reduces its size (ε > 0.15), but smaller voids continue to grow (ε = 0.20), until complete specimen failure. Concomitantly, the stress reaches a maximum level (at ε = 0.08) that progressively decreases until break. It should be remembered that both PP and PS are immiscible and no compatibilizer was used. The use of a compatibilizer should expectantly change the deformation mechanism [27, 60]. 17.3.4

Non-conventional Molding of PP Nanocomposites

Shear controlled orientation in injection molding (SCORIM) is a non-conventional injection molding techniques where high levels of shearing are applied to the molten polymer during the solidification phase [61]. This develops high level of molecular orientation. When processing clay-based polymer nanocomposites, these high shearings may be beneficial for achieving a high level of clay orientation and a better exfoliation and dispersion of the nanoclays, thus imparting an improved mechanical behavior of the moldings. 17.3.4.1

Materials and Sample Preparation

PP was reinforced with different amounts of incorporation of nanoclay (montmorillonite, MMT): 0, 0.5, 1, 3, 5 and 10 wt%. The PP and the nanoclays were mechanically blended in a tumbler mixer, and then directed injection molded in the form of a rectangular bar. High back pressure was used during the plasticating phase of the injection molding cycle in order to promote a better mixing. 17.3.4.2

Experimental Techniques

The molding microstructures were characterized by polarized light microscopy (PLM), scanning and transmission electron microscopy (SEM and TEM, respectively) and WAXS. Charpy impact tests were performed at room temperature (23 ◦ C). Fracture surfaces were observed by scanning electron microscopy (SEM). 17.3.4.3

Results and Discussion

Processing conditions and composition modify markedly the microstructure of molded PP nanocomposites. Figure 17.19 shows the effect of incorporation of MMT for fixed processing conditions (but expectantly for different thermomechanical environments due to the increase on melt viscosity by the incorporation of MMT). The addition of MMT affects the microstructure of the mouldings (e.g. skin–core structure, the development of semicrystalline morphologies). With increased content of MMT the microstructure becomes coarser and more multi-laminated. The MMT acts as a morphology director, affecting the structure development during processing. Figure 17.20 represents the SEM image of the external layer of the PP+5% nanoclay filled material system. In these images the flow direction is normal to the scanned surface, pointing outwards of the paper surface. A good dispersion of the MMT agglomerates in the polymer matrix was achieved. These agglomerates, with few

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PP + 0% MMT

PP + 0.5% MMT

PP + 1% MMT

PP + 3% MMT

Figure 17.19 Microstructure of the SCORIM moldings with different percentage of incorporation of nanoclay (MMT). The cuts are perpendicular to the bar length (or to the flow direction).

micrometers of length, have a platelet-like shape. They are all well aligned as response to the applied shear level during processing. SCORIM technique was capable of shaping the MMT agglomerates into platelets, of well dispersing them in the polymeric matrix, and of orienting them in the flow direction. Figure 17.21 shows the variation of the impact toughness of the moldings with increased content of MMT. In the same graph are also presented the correspondent 2D-WAXS patterns. The fracture energy is maximized for an incorporation of 5 wt% of MMT, with gains of more than 100% in the total absorbed energy. The addition of MMT does not change significantly the flexural modulus (not shown), but enhance greatly the impact toughness. Furthermore, low melt temperatures and high shearing times lead to the best mechanical performance. The mechanical properties of PP are determined by the incorporation of nanoclays and by the processing conditions, both being highly interrelated. The WASX patterns reveal that the skin layer of neat PP moulding shows a small level of crystalline phase orientation. This orientation increases with the incorporation of MMT. Figure 17.22 presents the equatorial I–2θ scan for different percentage of incorporation of MMT. Adding MMT induces the formation of β-phase PP (reflection at 2θ = 16.2◦ ), not present on the neat PP. This phase shows an improved toughness than the more common α-phase [62]. Simultaneously, the crystallinity index (not shown) decreases with the incorporation of MMT. The incorporation of MMT has several effects: (i) increases the level orientation of the crystalline phase; (ii) induces the formation of β-phase PP; and (iii) reduces the degree of crystallinity. The nanoclay acts as morphology directors, thus determining the mechanical response of the moldings.

Figure 17.20 Scanning electron microscopy image of PP filled with 5% of MMT. The flow direction points outwards the paper surface.

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Figure 17.21 Toughness of PP as function of % of incorporation of MMT nanoclay (SCORIM processing conditions: 12 strokes, stroke time of 3s and melt temperature of 280 ◦ C). The shown WAXS patterns are from the skin layer of the injection moulded impact bars.

17.3.5

Stretching of Nanoclay PET Nanocomposite

Polymer nanocomposites have received a lot of attention recently because of their improved performance, namely their enhanced mechanical behavior [63]. However, the mechanisms underlying these enhancements are still not well understood. The use of in situ structure-sensitive experimental techniques during deformation

Intensity (a.u.)

10% 5% 1% 0.5% 0%

10

12

14

16

18

20

22

24

26

28

2θ Figure 17.22 Equatorial Intensity-2θ curves for the skin layer of SCORIM molding with different % of incorporation of MMT (the WAXS patterns are those of Figure 17.21).

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studies are therefore of paramount relevance. Following this approach, small angle X-ray scattering (SAXS) investigations under synchrotron sources may give fundamental insights about deformation mechanisms at the nanoscale with adequate time-resolution. 17.3.5.1

Materials and Sample Preparation

R Different organo-modified layered silicate/montmorillonite were used: Nanofil 2 and 32 (referenced as NF2 and NF32, respectively) from Sud-Chemie. These nanoclays are functionalized with long chain hydrocarbon/benzyl groups. NF2 and N3F2 have average initial particle size of 8 and 30 μm, respectively (agglomerate dimensions), which are composed of platelets with typical dimensions of 100–500 nm × 1 nm (nanoclay). The PET nanocomposites with 0.3% of nanoparticles (and also neat PET for comparision purposes) were prepared via mechanical blending in a tumbler mixer and subsequent melt blending in a special asymmetric minimixer [64]. Compression molded plates were then prepared and rectangular tensile specimens were cut from them in a hydraulic press.

17.3.5.2

Experimental Techniques

A miniaturized uniaxial stretching machine was positioned perpendicular to the incident X-ray beam to perform tensile tests. Tensile specimens were stretched in situ at a constant velocity (5 mm/min) at HASYLAB, A2 soft condensed matter beam line, DESY, Hamburg, Germany, while acquiring two-dimensional SAXS patterns (30 s of accumulation time) 17.3.5.3

Results and Discussion

Figure 17.23 (top) shows the stress–strain curves of the PET nanocomposites, evidencing the different reinforcing nature of the nanoparticles, mainly in terms of stress and deformation levels. Both nanoclay types reinforced the PET, increasing the stress levels and the deformation capabilities. Figure 17.23 (bottom) also presents some SAXS patterns during stretching. The incorporation of the nanofillers influences the deformation mechanism of PET. This is dependent upon the initial size of the clays agglomerate (8 and 30 μm for NF2 and NF32, respectively). For the initial strains, ε = 0.22, the SAXS patterns are similar for neat and filled PET. For ε = 0.58, the SAXS patterns from neat and filled PET differ: the neat PET shows an intense horizontal streak corresponding to voids highly oriented in SD (between highly stretched fibrils that develop parallel to SD); whereas the PET nanocomposites show a cross pattern with an intense streak in the vertical direction (eventually due to external reflection of the interfaces of the crazes with the polymer) and a less intense streak in the horizontal direction (possibly due to elongated voids between fibrillis inside the craze) [66]. It may be suggested that the appearance of crazes is delayed (in terms of strain levels) for the PET nanocomposites. Furthermore, at that strain level the void size is slightly bigger for PET reinforced with NF2. At higher deformations, ε ≥ 0.95, and for all material systems, large voids are formed highly oriented in the SD, which increase in size with deformation. For the PET nanocomposites, the generated voids are smaller and with a narrower size distribution as compared with the neat PET. For the PET reinforced with NF32 (higher initial particle size), the formed voids are even smaller and have a narrower size distribution when compared to the NF2 nanoclay. PET-NF32 shows the lowest sustained stress level. The deformation mechanisms in polymer nanocomposites are still not well understood. The adoption of in situ structure-sensitive experimental techniques (mainly X-ray diffraction experiments) during deformation studies is therefore of paramount relevance for identifying the active deformation mechanisms and for establishing the relationships between the structure and mechanical performance.

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Figure 17.23 Stress–strain curve of PET nanocomposites (top) and correspondent SAXS patterns at different strain levels (bottom). Adapted from [65].

17.4 Concluding Remarks This chapter demonstrates that synchrotron WAXS and SAXS studies can be very useful in studying the relation between the structure and the mechanical properties in multiphasic polymer systems. The standard testing methods and software used for SAXS and WAXS data handling, however, show some limitations and need further development. In most of the instances, the WAXS and SAXS patterns are processed and interpreted, reflection by reflection, so as to extract indirectly the structural information (assuming a structural model). The latter can often be distorted or even damaged due to various reasons related either to the data collection or to the data treatment. The progress in the X-ray experiments during the last years has been tremendous and included development of new two-dimensional X-ray detectors, the use of high power X-ray microbeams, and the application of novel processing methodologies allowing for a direct transformation of

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the WAXS and SAXS 2D patterns into an image of the nanostructure. The calculation of the CDF, briefly discussed in this chapter, can provide structural information absolutely unavailable in other ways. With the advent of the nanotechnologies requiring a strict and rigorous control over the structures on the nanometer scale, this method can be of some importance for both industry and academia. The latest invention is the fast tomographic imaging method based on SAXS data from a scanning-microbeam experiment [67]. By means of this method, real-time X-ray experiments using mechanical testers for slow or fast load-cycling test can be incorporated into the synchrotron beamline. In such a way, fatigue and failure can be studied in polymer systems within reasonable intervals of time and the data related to microstructure variation inside the material.

Acknowledgements This work was supported by DESY and the European Commission under HASYLAB Projects DESYD-II-05-101 EC, DESY-D-II-07-011 EC and the FP6 contract RII3-CT-2004-506008 (IA-SFS). This work was also supported by FCT – Portuguese Foundation for Science and Technology through project POCTI/CTM/46940/2002 (MICROTEST) and POCI/CTM/57358/2004 (NANOFIBCO).

References 1. L. A. Utracki, in: Preface to Polymer Blends Handbook, Utracki L. A. (Ed.), 2002, Kluwer Academic Publishers, Dordrecht, The Netherlands. 2. L. S. Schadler, Polymer-based and Polymer-filled Nanocomposites in: Nanocomposite Science and Technology, Ajayan P.M., Schadler L.S. and Braun P.V. (Eds), 2003, Wiley-VCH Verlag GmbH Co. KGaA, Weinheim, p. 77. 3. A. N. Wilkinson, A.J. Ryan, Polymer Processing and Structure Development, 1999, Kluwer Academic Publishers, Dordrecht, the Netherlands. 4. G. Groeninckx and D. Dumpas, Plastic Deformation Mechanisms in Rubber and Rubber-Modified Thermoplastic Polymers: Molecular and Morphological Aspects. In: Structure and Properties of Multiphase Polymeric Materials, T. Araki, M. Shibayama, Q. Tran-Cong (Eds), 1998, CRC Press, Marcel Dekker Inc, New York, USA, p. 121. 5. D. Schwahn, Critical to Mean Field Crossover in Polymer Blends, Adv. Polym. Sci., 183:1–63 (2005). 6. K. Mortensen, Small-angle X-ray and Neutron Scattering Studies from Multiphase Polymers, Current Opinion in Solid State & Materials Science, 2: 653–660 (1997). 7. R. J. Roe, Methods of X-ray and Neutron Scattering in Polymer Science, Oxford University Press Inc., New York, 2000, pp. 90–128. 8. N. Stribeck, X-Ray Scattering of Soft Matter, Springer-Verlag, Berlin, Heidelberg 2007, pp. 42–61. 9. R. J. Davies, C. Riekel, J. A. Bennett, S. J. Eichhorn, R. J. Young, Probing the Internal Geometry of a Woven Composite During Deformation Using an X-ray Microdiffraction Imaging Technique, Appl. Phys. Lett. 91:044102 (2007). 10. C. L. Tucker III, P. Moldenaers, Microstructural evolution in polymer blends, Annual Review of Fluid Mechanics, 34: 177–210 (2002). 11. S. S. Ray, M. Okamoto, Polymer/layered silicate nanocomposites: a review from preparation to processing, Prog. Polym. Sci. 28 (2003) 1539–1641 12. S. C. Tjong, Structural and Mechanical Properties of Polymer Nanocomposites, Materials Science and Engineering R 53 (2006) 73–197 13. M. Evstatiev, S. Fakirov, Microfibrillar Reinforcement of Polymer Blends, Polymer 33:877 (1992) 14. S. Fakirov, M. Evstatiev, J.M. Schultz, Microfibrillar Reinforced Composite from Drawn Polyethylene Terephthalate/Polyamide 6 Blends, Polymer 34, 4669 (1993). 15. S. Fakirov, M. Evstatiev, S. Petrovich, Microfibrillar Reinforced Composites from Binary and Ternary Blends of Polyesters and Nylon 6, Macromolecules 26:5219 (1993).

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16. R. E. Newnham, J.R. Giniewicz, Non-mechanical Properties of Composites, in: Comprehensive Composite Materials, Vol. 1. Fiber Reinforcements and General Theory of Composites, Kelly A. and Zweben C. (Eds), Elsevier, Amsterdam, pp. 431–463 (2000). 17. S. Fakirov, M. Evstatiev, K. Friedrich, Nanostructured Polymer Composites from Polymer Blends: Morphology and Mechanical Properties, in: Handbook of Thermoplastic Polyesters, Fakirov S. (Ed.), Wiley-VCH, Weinheim, pp. 1093–1132 (2002). 18. S. Fakirov, M. Evstatiev, K. Friedrich, From Polymer Blends to Microfibrillar Reinforced Composites, in: Polymer Blends, Vol. 11, Formation and Preparation, Paul D.R (Ed.), Bucknall CB, J. Wiley and Sons, Inc., New York, p. 455–476 (2000). 19. A.A. Apostolov, M. Evstatiev, S. Fakirov, A. Kloczkowski, J.E. Mark, Structures and Mechanical Properties of Zone-Drawn-Zone-Annealed Blends of Cocrystallizing Poly(butylene terephthalate) and a Poly(ether ester, J Appl Polym Sci 59:1667 (1996). 20. M. Evstatiev, J. M. Schultz, G. Georgiev, S. Petrovich, S. Fakirov, K. Friedrich, In situ Polymer/Polymer Composites from Poly(ethylene terephthalate), Polyamide-6, and Polyamide-66 Blends, J Appl Polym Sci. 67:723 (1998). 21. D. Sapoundjieva, Z. Denchev, M. Evstatiev, S. Fakirov, N. Stribeck, M. Stamm, Transcrystallization with Reorientation in Drawn PET/PA12 Blend as Revealed by WAXS from Synchrotron Radiation, J Mater Sci 34:3063 (1999). 22. S. Fakirov, M. Evstatiev, Microfibrillar Reinforced Composites – New Materials from Polymer Blends, Adv Mat 6:395 (1994). 23. M. Evstatiev, S. Fakirov, G. Bechtold, K. Friedrich, Structure-Property Relationships of Injection- and CompressionMolded Microfibrillar-Reinforced PET/PA-6 Composites, Adv Polym Techn 19:249 (2000). 24. Z. Denchev, M.J. Oliveira, O.S Carneiro, Nanostructured Composites Based on Polyethylene-Polyamide Blends. I. Preparation and Mechanical Behavior, J Macromol Sci – Phys, B43:143 (2004). 25. S. Fakirov, H. Kamo, M. Estatiev K. Friedrich, Microfibrillar Reinforced Composites from PET/LDPE Blends: Morphology and Mechanical Properties, J Macromol Sci – Phys B43:775 (2004). 26. K. Friedrich, M. Evstatiev, S. Fakirov, O. Evstatiev, M. Ishii, M. Harras, Microfibrillar Reinforced Composites from PET/PP Blends: Processing, Morphology and Mechanical Properties, Comp Sci Techn 65: 107 (2005). 27. Z. Denchev, N. Dencheva, Transforming Polymer Blends into Composites: A Pathway towards Nanostructured Materials, Polym Int 57:11–22 (2008). 28. A. Monticciolo, P. Cassagnau, A. Michel, Fibrillar Morphology Development of PE/PBT Blends: Rheology and Solvent Permeability, Polym Eng Sci 38: 1882 (1998). 29. I. Pesneau, A. Ait-Kadi, M. Bousmina, P. Cassagnau, A. Michel, In-Situ Polymer-Polymer Composites, ANTEC 99, New York, May 2–6 1999; 2661. 30. M. Evstatiev, S. Fakirov,B. Krasteva, K. Friedrich, J.A. Covas, A.M. Cunha, Recycling of Poly(Ethylene Terephthalate) as Polymer-Polymer Composites, Polym Eng Sci 42:826 (2002). 31. D. Schmidt, D. Shah, E.P. Giannelis, New Advances in Polymer-Layered Silicate Nanocomposites, Current OpinionCurrent Opinion in Solid State and Materials Science 6, 2002, 205–212. 32. M. Okamoto, Recent Advances in Polymer/Layered Silicate Nanocomposites: An Overview from Science to Technology, Materials Science and Technology, 22(7), 756–779 (2006). 33. C. L Wu, M.Q. Zhang, M.Z. Rong, K. Friedrich, Tensile Performance Improvement of Low Nanoparticles Filled Polypropylene Composites, Compos. Sci. & Technol., 62:1327–1240 (2002) 34. H.R. Dennis, D.L. Hunter, D. Chang, S. Kim, J.L. White, J.W. Cho, D.R. Paul, Effect of Melt Processing Conditions on the Extent of Exfoliation in Organoclay-based Nanocomposites, Polymer, 42(23): 9513–9522 (2001). 35. Q. Zhang, Y. Wang, Q. Fu, Shear-induced Change of Exfoliation and. Orientation in Polypropylene/Montmorillonite Nanocomposites, J Polym Sci: Part B: Polym Phy, 41:1–10 (2003). 36. J.J. Hernandez, M.C. Garcıa-Gutierrez, A. Nogales, D.R. Rueda, A. Sanz, I. Sics, B.S. Hsiao, Z. Roslaniec, G. Broza, T.A. Ezquerra, Deformation Behaviour during Cold Drawing of Nanocomposites Based on Single Wall Carbon Nanotubes and Poly(ether ester) Copolymers, Polymer 48:3286–3293 (2007). 37. L. E. Alexander, X-Ray Diffraction Methods in Polymer Science, Wiley-Interscience, New York, 1969. 38. F. J. Balt´a-Calleja, C. G. Vonk, X-ray Scattering of Synthetic Polymers, Elsevier, Amsterdam, 1989. 39. W. Ruland, X-ray Determination of Crystallinity and Diffuse Disorder Scattering, Acta Crystallogr. 14:1180 (1961).

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40. Z. W. Wilchinsky, On Crystal Orientation in Polycrystalline Materials, J. Appl. Phys., 30, 792 (1959) 41. C. Santa Cruz, N. Stribeck, H. G. Zachmann, F. J. Balt´a-Calleja, Novel Aspects of the Structure of Poly(Ethylene Terephthalate) as Revealed by means of Small-Angle X-ray Scattering, Macromolecules 24:5980–5990 (1991). 42. G. Kortleve, C. G. Vonk, X-ray Small-Angle Scattering of Bulk Polyehtylene, Kolloid-Z. 225:124–131 (1968). 43. W. Ruland, Evaluation of Small-Angle Scattering of Anisotropic Lamellat Two-Phase Systems by means of Interface Distribution Functions, Colloid Polym Sci 256:932–936 (1978). ˇ 44. Z. Denchev Z, T .A. Ezquerra, TA, A. Nogales, I. Sics, Microstructural Characterization of Poly(Ethylene Naphthalene 2,6-Dicarbocylate) as Revealed by the Properties of Both Amorphous and Crystalline Phases, in: Handbook of Thermoplastic Polyesters: PET, PBT, PEN – homopolymers, copolymers, blends and composites, S. Fakirov (Ed.), 2002, Wiley-VCH, Weinheim, p. 483–546. 45. N. Dencheva, T. Nunes, M. J. Oliveira, Z. Denchev, Crystalline Structure of Polyamide 12 as Revealed by Solid-State 13 C NMR and Synchrotron WAXS and SAXS, J. Polym. Sci. Part B: Polym. Phys 43:3720–3733 (2005). 46. N. Dencheva, T. Nunes, M. J. Oliveira, Z. Denchev, Microfibrilar Composites based on Polyamide/Polyethylene Blends; 1. Structure Investigations in Oriented and Isotropic Polyamide 6, Polymer 46:887–901 (2005). 47. N. Dencheva, M. J. Oliveira, O. S. Carneiro, A.S. Pouzada, Z. Denchev, J. Appl. Polym. Sci. 115(5):2918-2932 (2010), John Wiley and Sons. 48. N. Preschilla, G. Sivalingam, A.S. Abdul Rasheed, S. Tyagi, A. Biswas, J. R. Bellare. Quantification of Organoclay Dispersion and Lamellar Morphology in Poly(Propylene)–Clay Nanocomposites with Small Angle X-ray Scattering, Polymer 49:4285–4297 (2008). 49. J. Bandyopadhyay, S. S. Ray, The Quantitative Analysis of Nano-clay Dispersion in Polymer Nanocomposites by Small Angle X-ray sScattering combined with Electron Microscopy, Polymer 51:1437–1449 (2010). 50. N. Dencheva, M. J. Oliveira, O. S. Carneiro, A.S. Pouzada, Z. Denchev, Preparation and Structure Development in Microfibrilar Composite Materials based on Polyethylene/Polyamide 6 Oriented Blends, J. Appl. Polym. Sci. 115(5): 2918–2932 (2010). 51. Nogales A, Hsiao BS, Somani RH, Srinivas S, Tsou AH, Balta-Calleja FJ, Ezquerra TA, Shear-induced Crystallization of Isotactic Polypropylene with Different Molecular Weight Distributions: In situ Small- and Wide-angle X-ray Scattering Studies, Polymer 42:5247–5256 (2001). 52. POLAR software version 2.7.1 is developed by Stonybrook Technology and Applied Research Inc. NY, USA 53. N. Stribeck, S. Fakirov, Three-Dimensional Chord Distribution Function SAXS Analysis of the Strained Domain Structure of a Poly(ether ester) Thermoplastic Elastomer, Macromolecules, 34:7758–7761 (2001). 54. V. Barbi, S. S. Funari, R. Gehrke, N. Scharnagl, N Stribeck, Nanostructure of Nafion Membrane Material as a Function of Mechanical Load Studied by SAXS, Polymer, 44:4853–4861 (2003). 55. Z. Denchev, N. Dencheva, S. S. Funari, M. Motovilin, T. Schubert, N. Stribeck, Nanostructure and Mechanical Properties Studied During Dynamical Straining of Microfibrillar Reinforced HDPE/PA Blends, Journal of Polymer Science: Part B: Polymer Physics, 48(3): 237–250 (2010). 56. J. P. Nunes, A.S. Pouzada, C.A. Bernardo, The Use of a Three-point Support Flexural Test to Predict the Stiffness of Anisotropic Composite Plates in Bending, Polym Testing 21:27–33 (2002). 57. M. van Duin, M. Aussems, R.J.M. Borggreve, Graft Formation and Chain Scission in Blends of Polyamide-6 and-6. 6 with Maleic Anhydride, J. Polym Sci Part A Polym Chem 36:179–188 (1998). 58. Z.-M. Li, B.-H. Xie, S. Yang, R. Huang, M.-B. Yang, Morphology-tensile Behavior Relationship in Injection Molded Poly(Ethylene Terephthalate)/Polyethylene and Polycarbonate/Polyethylene Blends (II) Part II Tensile Behaviour; J. Mat. Sci., 39:433– 443 (2004). 59. M. Fujiyama, Higher Order Structure of Injection-Moulded Polypropylene, in Polypropylene: Structure, Blends, and Composites, J. Karger-Kcosis (Ed.), Chapman and Hall, London, p.167–204 (1995). 60. S.A. Xu, S.C. Tjong, Deformation Mechanisms and Fracture Toughness of Polystyrene/High-Density Polyethylene Blends Compatibilized by Triblock Copolymer, J. Appl. Polym. Sci., 77, 2024–2033 (2000). 61. G. Kalay, R.A. Sousa, R.L. Reis, A.M. Cunha, M.J. Bevis, The Enhancement of the Mechanical Properties of a High-density Polyethylene, J. Appl. Polym. Sci., 73, 2473–2483 (1999). 62. S. C. Tjong, J.S. Shen and RK.Y. Li, Impact Fracture Toughness of β-form Polypropylene, Scripta Metallurgica et Materialia, 33(3), 503–508, (1995).

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63. S. N. Bhattacharya, M. R. Kamal, R. K. Gupta, in Polymeric Nanocomposites: Theory and Practice, Hanser, Munich, 2008. 64. O. Breuer, U. Sundararaj and R.W. Toogood, The Design and Performance of a New Miniature Mixer for Specialty Polymer Blends and Nanocomposites, Polymer Eng. Sci., 44(5), 868–879 (2004). 65. L. Todorov, A.D. Kouyumdzhiev, D.S. Teixeira, J.C. Viana, S.S. Funari, Evolution of SAXS Patterns during Stretching of PET Nanocomposites, Hasylab Annual Report 2007, part I, 1343–1342 (2007) http://hasyweb.desy.de/ science/annual_reports/2007_report/part1/contrib/46/22090.pdf 66. H.R. Brown, E.J. Kramer, Craze Microstructure from Small-angle X-ray Scattering (SAXS), J. Macromol. Sci., Part B, 19(3), 487–522 (1981). 67. Stribeck N, N¨ochel U, Fakirov S, Feldkamp J, Schroer C, Timmann A, Kuhlmann M, SAXS-Fiber ComputerTomography. Method Enhancement and Analysis of Microfibrillar-Reinforced Composite Precursors from PEBA and PET, Macromolecules 41:7637–7647 (2008).

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18 Characterization of Multiphase Polymer Systems by Neutron Scattering Max Wolff Department of Physics, Uppsala University, Uppsala, Sweden and Chair for Solid State Physics, Ruhr-University Bochum, Germany and Institut Laue-Langevin, Grenoble, France

18.1 Introduction A key issue within modern materials science is the development of materials based on soft matter, such as polymers and liquid crystals. The industrial use of polymers and related complex fluids comprises a wide range of applications, including the use of those special properties provided by the polymeric natures such as visco- and rubber-elasticity, adhesion and lubrication. One reason for this is that the properties can be tuned over a wide range to meet the requirements of application. In particular, multiphase polymer systems, including blends, block copolymers and semicrystalline polymers, are of high interest as they phase separate on mesoscopic length scales. This makes them interesting as templates for nanostructures [1–6]. The thermodynamics of polymer systems is dominated by the large number of degrees of freedom, weak interactions, delicate balance between enthalpy and entropy, large concentration fluctuations and macroscopic softness. The large molecular dimension of polymer systems markedly reduces the mixing entropy and provides the basis for self-organized structures. Neutron scattering is a particularly powerful technique for studying complex materials and has played a major role for the investigation of the basic properties of polymer systems. In recent years a variety of new auxiliary equipment has been invented, such as high pressure and shear.

18.2 Method of Neutron Scattering Several characteristic properties (Table 18.1) of the neutron makes scattering with this kind of probe valuable for the investigation of polymer systems in general and multiphase polymer systems in particular.

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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Handbook of Multiphase Polymer Systems Table 18.1 Properties of the neutron. Life time Mass Charge Spin Mag. moment

890 s 1.675 * 10−27 kg 0 1 2

μ/μn = −1.913

The fact that the neutron has a finite mass results in a dispersion E = 12 m n v 2 = 2mkn , where v and k denotes the velocity and wave vector of the neutron with mass m n , that is different from photons E = ck, where  denotes Plancks constant. For thermal neutrons (T = 293 K) this results in an energy of 25 meV and a wave length of 0.18 nm perfectly matching molecular excitations and inter-atomic distances. Of special importance for soft matter studies are the following aspects: 2 2

1. The wave length of thermal neutrons is comparable to inter-atomic distances. The coherence allows to address mesoscopic length scale as needed for phase separation in multiphase polymer systems. 2. The energy of thermal neutrons is low in comparison to photons of the same wave length and comparable to molecular and lattice excitations and allows to investigate diffusion. 3. The scattering power of light elements is huge. Moreover, hydrogen and deuterium have distinct scattering properties allowing contrast variation experiments. 4. The pentration power of neutrons for many engeneering materials is high. This makes them in particular well suited for the investigation of light elements in complex sample environments like, e.g. furnaces, cryostats or pressure and shear cells. 5. Neutrons have a magnetic moment and allow magnetic studies and spin echo experiments probing slow dynamics on the order of up to several 100 ns. 6. The low energy and weak interaction of neutrons with matter reduces sample damage. The time and length scales accessible to neutron scattering are illustrated in Figure 18.1 and compared to complementary techniques. 18.2.1

Scattering Experiment

In every scattering experiment the scattered intensity is detected under a defined angle. If additionally the energy of the incident and exiting waves are compared two relevant parameters can be extracted: 1. The energy transfer between the incident E and scattered E neutrons: ω = E  − E =

2 2 (k − k 2 ) 2m n

(18.1)

2. The momentum transfer between the incident k and scattered k neutrons: Q = k − k Variables are indicated in Figure 18.2.

(18.2)

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707

Length [nm] 5

4

10

3

10

102

10

101

100

10-1 100

10-16 Inelastic x-ray

10-14 Optical Ramansp.

10-2

scattering

10-12

Time [s]

Backscattering

&

10-6

10-8

Spectroscopy

10-10

Optical correlation

10-4

spectroscopy

10-2

PFG -

x-ray-

NMR

correlation-

10-12

100 10-5

10-14 10-16

spectroscopy

10-4

10-3

10-6

Spin-Echo

Farby- Perot-

10-8

10-4 Energy [eV]

10-10

TAX TOF

Optical Brillouin-

10-2

10-1

100

101

Scattering vector [nm-1]

Figure 18.1

Time and length scale accessible with neutron scattering.

Three scattering processes can be distinguished: 1. Elastic scattering: For elastic scattering the amount of the momentum of the incident k and exiting k neutrons is equal, only the direction changes. Equation 18.2 can then be rewritten by use of k = 2π : λ |Q| =

φ 4π sin λ 2

(18.3)

Detector Inelastic scattering E’2 >E k’2

Q

2

Elastic scattering E’=E k’ Qe Inelastic scattering E’<E 1

k’1 Θ Sample

Q

1

k

Figure 18.2 Scattering geometry: The vector Qe indicates the momentum transfer for elastic scattering, Q2 for the case of neutrons gaining energy during the scattering event and Q1 for the case of neutrons losing energy.

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k

k’

d hkl

s θ

θ

Figure 18.3 Bragg equation: Scattering at different lattice planes at distance dhkl results in a difference in path length s of the neutrons.

Where φ = 2 defines the scattering angle between the incident and the exiting wave vectors. In addition the coherence length of the wave (distance over which the phase is well defined and the wave can interfere) is preserved. Figure 18.3 depicts the scattering of an incident wave at lattice planes under an angle θ . The distance between lattice planes is noted by dhkl . In the detector, fixed under the same angle θ with respect to the lattice planes, intensity is only detected if the exiting waves interfere constructively. This happens if the difference in path s between waves scattered at two lattice planes meets exactly a multiple of the wave length: nλ = 2dhkl sin θ

(18.4)

2. Inelastic scattering: For inelastic scattering in addition to the change in momentum of the incident neutron beam the wave length of the neutrons changes. The coherence of the neutron beam is conserved. The change in energy is equal to the negative change in energy inside the sample. 3. Quasielastic scattering: For quasielastic scattering in addition to the direction of the momentum the coherence of the neutron beam is changed. The average wave length is preserved. The reduction in the coherence length is related to the dynamics and vibrations of the scattering particles on the time scale of the interaction and the resulting loss in phase information. A reduced coherence length means a less well defined energy of the neutron and results in a quasielastic broadening of the elastic line in the spectrum that can be correlated to the diffusive motion of the particles. In the scattering experiments a superposition of these effects is always recorded. By isotopic labeling the experimental conditions can be tuned in order to maximize the visibility of the desired scattering process. Depending on the physical question it can be decided whether the structure or the dynamics in a sample is investigated.

18.2.2

Born Approximation

In this section a short overview of the background of the scattering theory is given. An extensive background of the quantum mechanical description is given by Sakurai [7] or Sears [8]. The scattering potential can be related to the scattering cross section, which is experimentally accessible, by solving the

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Schroedinger equation: H | = E |

(18.5)

The Hamiltonoperator H contains the kinetic energy of the free neutron and the potential V(r) resulting from the interaction with the sample [9]: H = H0 + V (r)

(18.6)

For large distances (r → ∞), a weak scattering potential and by neglecting multiple scattering, the solution is given by an incident plane wave and an exiting spherical wave (first Born approximation): (r ) = (2π )

− 32

  eikr ikr e + f (Q) r

(18.7)

The function f (Q) = −

m 2π 2

 V (r)eiQr dr

(18.8)

is called scattering function and given by the Fourier transform of the scattering potential. In contrast to X-rays, neutrons interact with the core of the atoms. The size of the core (10−6 nm) is small in comparison to the wave length (0.1 nm) of thermal neutrons. Accordingly, the scattering process is isotrop and can be described by a single complex number b, the scattering length. This phenomenological constant has to be determined in an experiment. The real part describes the scattering power of the atoms and the imaginary part the absorption. The scattering potential for neutrons is defined as the Fermi-pseudo potential: V (r) =

2π 2 bi δ(r − R) m

(18.9)

The differential scattering cross section is the ratio of neutrons scattered into an infinitesimal angular element and the incident neutron flux given by the square of the scattering amplitude: ∂σ = | f (Q)|2 ∂

(18.10)

This quantity can be measured. The total scattering cross section σ S is defined as the average number of neutrons scattered normalized to the incident beam intensity:  σs =

 dσ =

| f (Q)|2 d = 4π |b|2

(18.11)

4π

The absorption cross section σa is defined as average number of neutrons absorbed normalized to the incident beam intensity and related to the imaginary part of the scattering potential: σa =

4π I m[b] k

(18.12)

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The sum of the scattering and the absorption cross section  gives the total scattering cross section (σt = σs + σa ) and is measured in barn 1 = barn = 10−24 cm2 The double differential scattering cross section, determined by the ratio of neutrons scattered with a certain energy transfer d E = dω into a certain angular element d and the incoming flux is given by:  m 2 k  ∂ 2σ = Pα |n  |V |n|2 δ (ω − ωα α ) ∂ ∂ω 2π 2 k α α

(18.13)

   |n = |k |α is the momentum |k and energy state |α of the neutron before and n  = k  α  after the scattering event. Pα is the occupation probability of the energy state |α before scattering. The scattering function is then defined as follows: S(Q, ω) =

4π k σs k 



∂ 2σ ∂ ∂ω

 (18.14)

This function S (Q, ω) is also called scattering law or dynamical structure factor and describes the structure and dynamics of the sample completely. 18.2.3

Elastic and Quasielastic Scattering

Generally the sample may consist of several different kinds of atoms each of which may have several isotopes. As shown in Figure 18.4 or Table 18.2, the scattering length of different cores may be distinct. Because of the magnetic moment of the neutron not only the composition of the core but also the full spin state of the neutron core system is important and the scattering length of different isotopes can vary. The average b of all spin states of one isotope is called coherent bcoh , and the standard deviation incoherent binc scattering length: bcoh = b

1 binc = b2  − b2 2

(18.15) (18.16)

From the values of b the coherent σ coh and incoherent σ inc scattering cross sections can be calculated: σcoh = 4π b2   σinc = 4π b2  − b2

H

(incoherent )

Figure 18.4

D

C

O

(18.17) (18.18)

Al

Si

Negative coherent scattering l. Positive coherent scattering l.

Scattering length: Scattering length densities for different atoms.

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Table 18.2 Scattering length of selected atomic cores extracted from Sears [8] for thermal neutrons v = 2200 m/s. Z indicates the atomic number and Am the atomic mass. Spin Z H 1 H 2 H 3 H He

B

Am

[]

natural occurance

1 2 3

1/2 1 1/2

99,985 0,015 (12,32A)

3

1/2

0,00014

4

0

99,9998

10

3

20

11

−3/2

80

12

0

98,9

16 27

0 5/2

99,762 100

28 31

0 1/2

92,23 100

Scattering length [fm] Scattering cross section [barn]

1

2

5

C

6

O

8

Al Si

13 14

P Cd

15 48

Coh.

Inc.

Coh.

Inc.

Scat.

Abs.

−3,74 −3,74 6,67 4,79 3,26 5,74 −1,48i 3,26 5,3 −0,21i −0,1 −1,07i 6,65 6,65 6,65 5,8 5,8 3,45 4,15 4,11 5,13 4,87 −0,7i

− 25,3 4,04 −1,04 − −2,5 +2,57i 0 −

1,7586 1,7599 5,597 3,07 1,34 4,42

79,90 79,91 2,04 0,0 0,0 1,2

81,66 81,67 7,64 3,07 1,34 5,6

0,3326 0,3326 0,00051 0 0,00747 5333

1,34 3,54

0 1,70

1,34 5,24

0 767

−4,7 +1,2i −1,3 − 0 − 0 0,26 − 0 0,2

0,14

3,0

3,1

3837

5,56 5,554 5,563 4,235 4,235 1,495 2,163 2,119 3,307 3,3

0,22 0,001 0 0,000 0 0,009 0,015 0 0,006 2,4

5,78 5,555 5,563 4,235 4,235 1,504 2,178 0,988 3,313 5,7

0,0055 0,00350 0,00353 0,00019 0,00010 0,231 0,171 0,54 0,172 2520

The differential scattering cross section splits into two parts ∂ 2σ = ∂ ∂ω



∂ 2σ ∂ ∂ω



 + inc

∂ 2σ ∂ ∂ω

 (18.19) coh

and the coherent and incoherent scattering functions are defined as follows: 

 ∂ 2σ ∂ ∂ω coh  2  ∂ σ 4π k Sinc (Q, ω) =  σinc k ∂ ∂ω inc

4π k Scoh (Q, ω) = σcoh k 

(18.20) (18.21)

To catch the meaning of the scattering function the fourier transform is defined as K(r, t) [10]: K (r, t) =

1 (2π )3



e−i(Qr−ωt) S(Q, ω)dQdω

(18.22)

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K (r, t) has the form of a correlation function: 1 K (r, t) = N





dr



 



δ(r − Rj (0))δ(r + r + Ri (t))

(18.23)

ij

Formally the single particle correlation function can be extracted from Eq. (18.22): 1 K s (r, t) = N



dr



 δ(r − Rj (0))δ(r + r + Ri (t)).

(18.24)

i= j

The fourier transform of the correlation function K in space and time is per definition Scoh and that of Ks , the single particle correlation function, is Sinc :  Scoh (Q, ω) =

(18.25)

K s (r, t)ei(Qr−ωt) drdt

(18.26)



Sinc (Q, ω) =

18.2.3.1

K (r, t)ei(Qr−ωt) drdt

Coherent Scattering Function

Integrating the coherent scattering function over energy yields the structure factor:  Scoh (Q) =

Scoh (Q, ω)dω

(18.27)

This represents a snapshot of the particle pair correlation function [11]. Scoh (Q) describes the Bragg reflections, liquid structure factor and phonons. It contains information about the ordering of atoms in space and the temporal evolution of this. 18.2.3.2

Incoherent Scattering Function

Integrating the incoherent scattering function over all energies gives:  Sinc (Q) =

Sinc (Q, ω)dω = 1

(18.28)

Sinc (Q) shows that particles are there, without information about correlations between them. Sinc (Q, ω) makes the self interference of the single particle visible and accordingly provides information about the self motion of particles. At time t = 0 the probability to find a certain particle at the position r can be assumed 1 (limt→0 K s (r, t) = δ(r)). For free diffusion this probability decreases to 0 (limt→∞ K s (r, t) = 0) with time. For spacial constraints, limiting the motion to a certain volume, the probability will decrease to a certain discrete value. The self correlation function can then be divided into two parts [12]: K s (r, t) = K s (r) + K s (r, t)

(18.29)

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Here Ks (r) is the probability to find the particle in distance r from the starting point after infinite time. This function does not depend on time and the Fourier transform is elastic in energy:  F(Q) =

K s (r)eiQr drδ(ω)

(18.30)

This function is directly related to the geometry of the diffusive path and is called elastic incoherent structure factor (EISF). The large difference in the scattering length of hydrogen and deuterium (see Table 18.2) allows, especially for polymers, to optimize samples with respect to the physical question in order to get predominant coherent or incoherent scattering. For a random jump diffusion, e.g. a Gaussian distribution of the intermediate scattering function can be assumed: I (Q, t) = e−Dt Q

2

(18.31)

This means for the scattering function: S(Q, ω) =

D Q2 1 π ω2 + (D Q 2 )2

(18.32)

This is a Lorentz function with the full width at half maximum (FWHM) 2: 2 = 2D Q 2

(18.33)

where D symbolizes the diffusion constant of the particle. 18.2.4

Scattering at Small Momentum Transfer

Multiphase polymer systems are characterized by correlations on mesoscopic length scales. Accordingly mostly scattering methods at low momentum transfer are used to gather information about these systems. The most important techniques in this context are small angle neutron scattering and, for additional surface sensitivity, reflectometry or gracing incidence scattering. 18.2.4.1

Neutron Small Angle Scattering

Conventional small angle neutron scattering (SANS) works in transmission whereas neutron reflectometry (NR) uses a reflection geometry. Accordingly, SANS probes the bulk properties of a sample whereas in NR the surface is studied. For both techniques the radius of the Ewald sphere or the amount of the incident wave vector, for example for a wavelength of 0.5 nm (k = 12.5 nm−1 ), is huge in comparison to the reciprocal lattice vectors, typically smaller than 1 nm−1 , under investigation. This allows the excitation of several Bragg reflections at the same time. It is important to note that for SANS the Born-approximation, assuming one single scattering event, holds, whereas neutron reflectivity is a dynamical or multiple scattering effect. In contrast to the scattering theory described above where the Schroedinger equation was solved for a single particle potential for small momentum transfers the details of the atoms become invisble for the scattered neutron. The phase shift between the scattered waves from neighboring atoms is small and smearing of the phase due to the finite coherence of the beam does not allow atomic resolution. Accordingly, over a certain (the coherence volume of the neutron beam) area averaged potential is probed. For example V(r) for atoms

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enclosed in a defined volume has the following form:  V (r) =

V0 0

inside outside

(18.34)

   2 where V0 = 2πmh bcoh ρn  and ρn (r ) = j δ r − r j is the number density of scatterers. In a small angle scattering experiment the Fourier transform of the density distribution is detected as differential scattering cross section:



2



∂σ iQr 3

= bcoh ρn (r)e d r

∂

(18.35)

This allows to determine the form factor of mesoscopic objects as well as the correlations between them. Neutron scattering experiments are particularly useful for studies of the detailed structures of colloidal and polymeric systems. This is done by systematic studies of systems in which the contrast is varied by exchanging hydrogen with deuterium in specific places. In a solution study, this is most easily done by changing the composition of the solvent in terms of normal and deuterated material. While the effect of changing the contrast of an essential two-phase system is only to change the prefactor in front of the scattering function, more complex structures with varying scattering length density will give rise to significant different forms of the scattering function. 18.2.4.2

Neutron Reflectometry

Now let us assume that the interface of V is in the z = 0-plane (Figure 18.5) and k in the y = 0-plane. The general solution of Eq. (18.5) is then:  ψ(r) =



inside z > 0

AeiKr ae

ikr

 ik r

+a e

outside z < 0

(18.36)

If the energy of the neutron is smaller than the potential barrier E 0 < V0 the neutrons are totally reflected. 2 2 For larger energies they are partly reflected and partly transmitted. If the amount of the component h2mk perpendicular to the interface is smaller than V 0 the wave cannot penetrate into the material and is totally

E V E E‘ V E o
Vo

Vacuum

Mater

x

Figure 18.5 Total external reflection: For slow neutrons matter is as a potential barrier V0 . Inside the following relation holds: E = E0 − V0 . If E0 < V0 the neutrons are total externally reflected.

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Characterization of Multiphase Polymer Systems by Neutron Scattering

reflected externally. Snell’s law n =

cos θ cos 

715

gives the angle of total external reflection: cos θc = n

(18.37)

Generally this angle θc is small (< 1◦ ) for thermal neutrons and a Taylor expansion of the sin gives:  sin θC ≈ θC =

1  v0 ρn bcoh 2 =λ E π

(18.38)

The amplitude of the refracted intensity can be calculated: A = aei(ωt−kx−k(θ

2

2

1

λ − 2π ρn b) 2 z)

(18.39)

For a complex scattering length b = b + ib the refracted beam is partly absorbed. The average penetration depth (depth, at which the beam is attenuated to 1/e) close to total external reflection is [13]: 1 λ 1 2 z 1/e = √ [(θ 2 − θc2 )2 + β 2 ] 2 − (θ 2 − θc2 ) 4 2π

(18.40)

n σabs describes the absorption. Figure 18.6 depicts the penetration depth plotted for different β. For β = λρ4π strong absorption the penetration depth even above the critical edge is small. For small absorption the penetration depth jumps to relatively large values at the critical edge. For angles larger than the critical angle this expression rapidly approaches the known Lambert-Beer expression [13]:

ρn σabs I (z 1/e ) = e−ρn σabs d1/e = e− θ −θc z1/e I (0)

(18.41)

Here d1/e denotes the path length of the neutrons inside the material along the direction of k, where the intensity drops to 1/e. As the incoherent scattering is independent of the angle and the phase information is lost, it can be treated as additional absorption in a diffraction experiment.

Penetration depth [nm]

 = σabs + σinc σabs

6

(18.42)

-10

10

absorption β = 10 -8 absorption β = 10 -6 absorption β = 10

5

10

4

10

3

10

2

10

1

10

0.0

0.5

1.0

α/αc

1.5

Figure 18.6 Penetration depth: For weakly absorbing materials the penetration depth increases jump-like at the critical edge.

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For the reflected intensity one obtains similar to optics the Fresnel reflectivity: R F = kc4 /Q 4z

(18.43)

The reflected intensity for different interfaces is summarized in Figure 18.7. In the Born approximation the reflected intensity can be described as follows:





1 ∂ V (z) i Q z z

2

R(Q z ) ∼ R F (Q z )

e dz δ(Q x )δ(Q y ) v(z) ∂z

(18.44)

For a perfectly flat interface this results in the Fresnel reflectivity (RF ). For a Gaussian distribution of the scattering length along the z-direction one obtains: bρn (z) = er f (z) = √

1 2π σ



z

−u 2

e− 2σ 2 du

(18.45)

0

where σ denotes the width of the Gaussian distribution. For the reflected intensity this results in the following expression: R(Q z ) = R F (Q z )e−Q z σ 2

2

(18.46)

The reflectivity for a rough interface drops more steeply than Q 4z (Figure 18.7). For the above assumptions there is no difference whether the smearing in scattering length results from real roughness or is due to interdiffusion (Figure 18.8). Information on the in-plane correlations is only accessible if the non-specular or diffuse scattering is analyzed [14].

18.3 Experimental Techniques In the following section, some technical details of neutron scattering are discussed. First the production, distribution and detection of neutrons at large-scale facilities and then the different kinds of instruments is described. 18.3.1

Production and Detection of Neutrons

In order to allow scattering experiments neutrons have to be provided and detected. 18.3.1.1

Neutron Sources

Two types of neutron sources exist: the reactor and the spallation. In the reactor core fast neutrons, i.e. neutron with energy En of the order of 1 MeV, are produced in a fission chain reaction. Usually enriched uranium 235 U is used for fuel. Almost all reactor facilities act as continuous sources, providing a constant neutron flux. In the reactor the core is surrounded by a moderator, typically made of D2 O, H2 O or graphite. By interaction with the moderator material, the neutrons are thermalized to the energy given by the temperature of the moderator, e.g. 50 ◦ C. The resulting Maxwell Boltzmann distributed

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Characterization of Multiphase Polymer Systems by Neutron Scattering 3

SLD [10 ]

2

Intensität

Silicon interface

10 10

1 0

10 10 10 10 10

-1 -20

-10

0

10

20

0.00

0.05

0.15

0.20

0.15

0.20

0.15

0.20

10 10

2

Intensität

SLD [10 ]

3

Strong absorption

0.10

Q [Angstr.-1 ]

Depth [Angstr.]

1 0

10 10 10 10 10

-1 -20

-10

0

10

20

0.00

0.05

0.10

Q [Angstr. -1]

Depth [Angstr.] 3 2

Intensität

Roughness

SLD [10 ]

10

1 0

10 10 10 10

-1 -20

-10

0

10

0.00

20

0.05

3

10

2

10

Intensität

SLD [10 ]

Layer

0.10

Q [Angstr.-1 ]

Depth [Angstr.]

1 0

Kiessig oscillations

10 10 10

-1 -50

0

50

100

150

200

250

0.00

0.05

5

10

4

10

3

10

2 1 0

-50

Kiessig oscillations

10 10 10

0

50

100

150

200

250

0.00

0.05

Depth [Angstr.]

0.10

0.15

0.20

Q [Angstr. -1] 10

4

10

3

10

Intensität

5

2 1 0

Bragg reflections

10 10 10 10

-1 0

1000

2000

Depth [Angstr.]

Figure 18.7

0.20

10

-1

Multilayer

0.15

Q [Angstr. ]

Intensität

Layer

0.10

-1

Depth [Angstr.]

SLD [10 ]

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Scattering length density profile (left) and resulting reflectivities (right).

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Figure 18.8

Roughness and interdiffusion.

thermal neutron flux has a maximum for neutron energies e.g. of E n = k B T = 30 meV, corresponding to a neutron velocity of  vn =

2k B T mn

(18.47)

of about 3000 m/sec, and a neutron wavelength of λ = 0.14 nm. At a spallation source a proton beam is sent on a liquid target, e.g. mercury. The high energy proton pulses excite the core of the target and result in the emission of neutrons. This process is more efficient than fission as 20–30 neutrons are emitted from one core whereas only three are emitted during fission, one of which is needed for the chain reaction. Moreover, the pulsed structure of the neutron beam can be tuned by changing the proton pulse length and repetition rate and directly be exploited for the instrumental needs and especially spectroscopic measurements profit from a huge gain in intensity. Neutrons are then delivered via neutron guides to the instruments. Such neutron guides are glass channels coated with Ni or super-mirrors reflecting neutrons with wavelength above a certain critical value, typically ˚ depending on the Ni-isotope that is used, or the design of the super-mirror. more than 3 A, 18.3.1.2

Monochromator

Neutrons with a particular wavelength can be extracted from the white spectrum of the neutron source by monochromators. In pulsed facilities, the obvious way is to use time of flight, even though additional monochromatization may be useful. For continuous sources, the monochromatization is essential. Three basic techniques are used, given by the dual properties of the neutron: respective wave and particle properties. Using the wave nature of the neutrons, a particular wavelength can be selected via Bragg scattering from a crystal with suitable crystal lattice spacing, d. A high reflectivity is needed to match the general resolution of the instrument, a relatively high mosaic spread is favored, thereby fulfilling the Bragg condition for neutron wavelength within some limited bandwidth. Typically, pyrolitic graphite is used as crystal monochromators. Alternatively, one can use the particle nature of the neutrons to choose the appropriate neutron velocity, and thereby given wavelength. The velocity of neutrons of 0.6 nm, as an example, is 600 m/sec. This relatively low speed allows the use of mechanical selectors made up of rotating plates of neutron-absorbing materials. Those neutrons with too low speed will be absorbed by the selector from the back, while those with too high speed will hit the selector plate in front. Mechanical selectors are usually made as slightly screwed turbine-like devices. The screw angle and the distance between the plates define the wavelength spread, while these parameters and the speed of rotation give the nominal wavelength. By changing the axis of rotation relative to the beam direction an additional variation of wavelength spread is possible. Another possibility, exploiting the particle character of the neutron, is the use of chopper discs in order to provide a pulsed and polychromatic or monochromatic beam. The principle is the same as for the velocity selector. The incident neutron passes a couple of chopper discs made of absorbing material with a hole to

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719

transmit the neutron beam. The first chopper cuts the white neutron beam into pieces wheres the second one selects a certain wave length. To provide frame overlap of subsequent pulses and suppress harmonics additional chopper discs are necessary. 18.3.1.3

Neutron Detector

Depending on the need of the experiment, instruments are either equipped with a 0-dimesional point detector or an area detector. The detector is usually of the single or multiwire type, with 3 He as the active gas for neutron detection. 3 He captures neutrons through the process n + 3 He → 3 H + 1 H + 0.76 MeV

(18.48)

producing tritium, (3 H), and protons, (1 H). Alternative gases or materials based on either 10 B of 6 Li can be used for neutron detection. In the gas detector there is in addition an ionizable gas of, for example, mixed argon and methane. The positive protons and tritium nuclei produce a trace of ionized Ar-gas that is accelerated towards wire planes, thereby making a charged pulse that is analyzed in the electronic system to provide the coordinate of the scattered neutron. The gas detector has a limited capability in terms of neutron flux and spatial sensitivity. For reflectometer detectors with a diameter of the order of 30 cm, the typical spatial resolutions used are within the 2 mm range. Dead-time problems start to become significant for count rates of the order of 105 to 106 counts/sec when distributed over the whole detector, whereas count rates within a few pixels normally should be limited to a few thousands. 18.3.2 18.3.2.1

Instrumentation Small Angle Neutron Scattering

The SANS instrument is in essence rather simple. Figure 18.9 shows the schematic layout of the D22 small angle scattering maschine at the Institut Laue-Langevin (Grenoble, France) [15]. Neutrons are monochromatized by a velocity selector. The neutron beam direction is defined by the collimator, which in principle can be given by just two pinholes, one near the source point, and one near the sample. In practice, the SANS facility, as D22, is often equipped with a variable length of the collimator as well as variable pinhole sizes. The insert in Figure 18.9 shows the setup by which either a piece of neutron guide or a collimator (polarizing

Velocity selector

Collimation

Neutron guides

Diaphragm

Sample

Detector Another position

Evacuated tube (20m)

Figure 18.9 Small angle scattering instrument: Schematic layout of the small angle scattering machine D22 at the ILL [15].

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Diaphragm Polariser

Diaphragm Sample Beam stop

Evacuated

Shutter flight tube

Be filter

Figure 18.10

Monitor

Spin flipper

Analyser

Diaphragm

Detector

Reflectometer: Schematic layout of the reflectometer ADAM at the ILL [15].

and nonpolarizing) can be rotated into the beam. The effective length is varied by inserting sections of neutron guides between the source (or the monochromator) and the sample. The diaphragm in front of the sample is for background reduction. The position sensitive detector can be moved forth and back inside a huge evacuated detector tank. As an example, a 15 m collimator with 10 mm diameter pinhole at the entrance, and 5 mm diameter pinhole near the sample corresponds to a beam divergence of 0.058◦ . For an equal detector distance and a wavelength of 1 nm this would result in a minimum correlation length of 5 μm that could be resolved. For larger scattering angles (larger than .1◦ ) the resolution is determined by the wavelength spread of usually 10% of the incident beam. 18.3.2.2

Reflectometer

Figure 18.10 shows the setup of the angle dispersive reflectometer ADAM at the Institut Laue-Langevin (ILL, Grenoble, France) [16]. The principle is the same as for a diffractometer, albeit optimized for scattering under very shallow angles. This implies a higher demand on the divergence in the scattering plane. To compensate for neutron flux the resolution in the perpendicular direction can be relaxed. The neutron beam is monochromatized by a vertically focusing HOPG monochromator and the higher order wave lengths are suppressed by a transmission Be-filter. The beam is then collimated by two variable vertical slits defining the resolution. A polarizer allows for optional beam polarization. Neutrons are finally detected at a positionsensitive detector. In contrast to SANS, at a reflectometer the resolution is usually given by the divergence and thus setting of the two collimating diaphragms, of the incident beam rather than by the wavelength spread. Another possibility for neutron reflectometry experiments is to use the different velocities of the neutrons and time of flight of a pulsed neutron beam. The different wave length result similar to the different angles for the angle dispersive version in different values for the momentum transfer while taking data at a fixed angle of incidence and positon of the detector. This has the advantage that part of the reflectivity is measured in one go. On the other hand, data have to be normalized to the incident wave length spectrum. 18.3.2.3

Time of Flight

Figure 18.11 depicts the setup of a time of flight spectrometer. In the primary spectrometer a pulsed neutron beam is generated, which is analyzed in the secondary spectrometer after scattering at the sample position. The primary spectrometer consists of several chopper disks spinning around a horizontal axis. The first chopper produces a white pulsed neutron beam. The second chopper disk monochromatizes the beam and

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Detectors

4m

Beam stop

Choppers

Focusing neutron guide

Choppers

Monitor Sample

Figure 18.11

Radial collimator

1.2 m

8m

Time of flight spectrometer: Schematic layout of the instrument IN5 at the ILL [15].

thus selects a certain neutron energy or wave length: λ=

h h = L mv mτ

(18.49)

where L is the distance between the chopper disks and τ the flight time of the neutrons for this distance. The phase between the two disks is defined modulo 2π and a third chopper with the same rotational speed is needed to filter for higher harmonics. A fourth chopper rotating at a lower frequency is filtering intermediate pulses and assures a certain time interval between the pulses long enough to get a complete time analysis of the different pulses and prevent overlapping of subsequent ones. The secondary spectrometer consists of the sample table and the detectors covering a large angular volume. From the difference of the flight time of the inelastic neutrons in comparison to the flight time of the elastically scattered neutrons the energy transfer can be calculated: mn 2 L E = 2 PD



1 1 − 2 2 τel τinel

 (18.50)

Here LPD is the distance between the sample and the detector, τinel the flight time of the inelastically and τ el that of the elastically scattered neutrons. 18.3.2.4

Triple Axis – Back Scattering

On a triple axes spectrometer (Figure 18.12) a certain wave length is selected with a monochromator out of the white spectrum in the neutron guide. The energy analysis after the sample is done with a rotatable crystal analyzing the outgoing wave length under a certain angle or Q value. On a backscattering instrument both monochromator and analyzer are used in backscattering condition. The resolution of these kinds of instruments is essentially given by the monochromator, the analyzer crystal and the angular divergence. At both crystals the neutrons are Bragg scattered. The energy uncertainty in the

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Shutter

Neutron Guide

Be filter Focusing guide 1 Monitor 1 Collimator or Bender Sample

Diaphragms

Collimator

Beam stop

Monitor 2

Collimator Detector

Analyser

Figure 18.12

Triple axis spectrometer: Schematic layout of the triple axis instrument IN14 at the ILL [15].

Bragg equation can be determined 2λ 2d E = = 2 cot θ θ + E λ d

(18.51)

The second term is given by the quality of the crystal, whereas the first one vanishes for θ = 90◦ . The quality of the crystal is given by deviations from the mean lattice parameter and the mosaic spread. Practically, these limits are never reached and the resolution of a backscattering instrument is always given by the angular divergence. Figure 18.13 depicts the setup of the backscattering instrument IN16 (ILL). A deflector is extracting neutrons from the guide. Then a background chopper cuts the continous neutron beam into pulses. A second chopper prevents neutrons directly hitting the detector and frame overlap of different pulses. A second deflector directs neutrons towards the monochromator. As the monochromator is used in backscattering condition this deflector is a quarter circle disk. The monochromator generates a neutron beam with a very well-defined wave length. To measure different energy transfers at the sample position the monochromator

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Analysers

Second deflector

Helium flight box Collimator Detector

Chopper Sample Neutron guide

Figure 18.13

Be filter

Moving Monochromator First deflector Beam stop

Focusing guide

Back scattering instrument: Schematic layout of the instrument IN16 at the ILL [15].

is moved along the neutrons, flight direction. As the velocity of the monochromator and the flight time of the neutrons to the detector are known the energy can be determined. The neutrons scattered at the sample are analyzed by huge analyzer plates covering a large solid angle. These plates only reflect a very well-defined wave length and only neutrons can be detected that have exactly the same energy shift but in opposite direction at the sample and the monochromator, respectively. Spacial resolution of the detectors allows to register the angular and energy dependence of the scattering process. 18.3.2.5

Neutron Spin Echo

Figure 18.14 depicts the schematic layout of the IN15 Neutron Spin Echo (NSE) instrument at the ILL. The idea of this technique is to exploit the spin of the neutrons in order to achieve a very high energy resolution without the need of a well collimated or highly monochromatic neutron beam. This is realized by use of a polarized neutron beam passing two equivalent magnetic regions (precession solenoids) with field directions perpendicular to the neutron spin and along opposite directions before and after the scattering process. A neutron with spin 1/2 will precess in a perpendicular magnetic field. The total precession angle is given by: φ =

γ Bl v

(18.52)

where γ is the gyromagnetic ratio of the neutron, B the magnetic field, l the path length of the neutron inside the magnetic field and v the velocity of the neutron. For no energy transfer full polarization is recovered at the analyzer. As the precession angle is propotional to the field integral a long distance or high

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Normal version with neutron guide Selector

Polarizer

π flipper coil 2

π flipper coil 2 Small echo

Sample

Sample field coil

Supermirror analyser

Area detector

Beam stop

Precession Precession solenoid 1 solenoid Small echo

π flipper coil

Or four chopper system for TOF mode

Polarizer

Neutron focusing mirror

Mirror version with neutron focusing

Figure 18.14

Neutron spin echo instrument: Schematic layout of the spectrometer IN15 at the ILL [15].

magnetic field is desirable. In practice, a long distance has the disadvantage of intensity loss due to the beam divergence. Accordingly, on spin echo instruments high fields are provided in the coils. As the polarization recovered independent of the wavelength a rather relaxed wavelength resolution can be used on a spin echo instrument if no good Q resolution is needed. Accordingly, for most experiments on IN15 a velocity selector providing 15% wave length spread is used. Neutrons are polarized by a supermirror polarizer. As it is difficult to change the direction of the magnetic field by 180◦ as this would imply a zero field region that may depolarize the beam, a π/2 flipper changes the spin direction of the neutron with respect to the precession fields. The resolution of a spin echo instrument is then mainly defined by the homogeneity of the magnetic field. For an inhomogeneous magnetic field neutrons on different trajectories will experience a different phase shift which will be measured as depolarization similar to a transfer of energy during the scattering process.

18.3.3

Grazing Incidence Small Angle Scattering

To address in-plane ordering for many reflectivity studies a position-sensitive detector (PSD) is used for data collection to catch the off-specular scattering. For a standard measurement the optimum resolution is used only in the scattering plane. To increase the intensity a divergent incoming neutron beam is focused on the sample in the perpendicular direction (out of the scattering plane). Improving the resolution out of the scattering

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725

Position–Sensitive Detector z y

k|

k Silicon

x

Sample

Figure 18.15 Scattering geometry for GISANS experiments. Reprinted from [18]. Copyright (2004) with permission from American Physical Society.

plane allows to collect SANS information simultanously. Figure 18.15 shows the scattering geometry for Grazing Incidence Small Angle Neutron Scattering (GISANS) studies. Neutrons from the monochromator are reflected at an interface and registered by a two dimensional position-sensitive detector. The z-direction is defined by the normal to the interface, the x-direction is in the scattering and interface planes and the y-direction is perpendicular to x and z. To elucidate the interference between SANS and NR, Figure 18.16 shows a comparison between the GISANS scattering pattern (left panel) taken at an incident angle of 0.3◦ , for a sample of a 20% solution of F127 in deuterated water at 298 K, and a map where the reflectivity data, including the off-specular scattering, is plotted over Qx and Qz (right panel) for similar experimental conditions [17]. Aqueous solutions on F127 are known to form micelles that crystallize under certain conditions. This sample system is later described in more detail. The only difference between the measurements was that the resolution in Qy -direction was increased for the GISANS measurement. The black line, plotted in both panels of Figure 18.16, indicates where the two hyper surfaces intersect. For the relaxed resolution the double peak showing up in the GISANS data at Qy = 0 is detected in the scattering plane. Low

Intensity

High

Qz{10–2 Å–1}

7.5

7.5

5

5

2.5

0 –7.5 –5

Qz[10–2 Å–1]

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–2.5 0 2.5 5 Qy[10–2 Å–1]

7.5 –7.5 –5 –2.5 0 2.5 Qx[104 Å–1]

5

0 7.5

Figure 18.16 Comparison of GISANS (left panel) and reflectivity (right panel) data. Reprinted from [115]. Copyright (2005) with permission from European Physics Journal.

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10

Length scale [1/ ]

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10

dx dy, d

6

10

z

4

10

2

10

0

10

-4

10

-3

10

-2

10

-1

10

0

10

1

10

Exit angle [°]

Figure 18.17 Length scales probed for the different directions x, y and z. Reprinted from [115]. Copyright (2005) with permission from European Physics Journal.

The three components of Q, with respect to the coordinate system defined in Figure 18.15, can be written down as follows: 2π (cos α − cos α  cos α⊥ ) λ 2π sin α⊥ Qy = λ 2π (sin α + sin α  ), Qz = λ

Qx =

(18.53) (18.54) (18.55)

where α and α  are the incident and exit beam angles out of the interface plane, respectively, and α ⊥ is the scattering angle in the plane of the sample surface relative to the x–z plane. It turns out that Qz and Qy show a sin dependence and increase for small angles linearly with the scattering angle (α + α  ) and α ⊥ , respectively. Apart from that, Qx is calculated from the difference between the cos of α and α  multiplied with α ⊥ . For small angles the cos shows a α 2 dependence and is nearly one meaning a much smaller component of Q in normalized to the wave length probed in the three x-direction. Figure 18.17 shows the length scale di = 2π Qi directions calculated for α = 0. It turns out that for angles α  and α ⊥ < 1◦ the length scale under investigation in x-direction dx is at least two orders of magnitude bigger than in y and z-direction (dy and dz ). 18.3.4

Comparison of SANS and GISANS for Crystalline Systems

Figure 18.18 shows the intensity taken with a sample of F127 solved in deuterated water plotted in a map over Qx and Qy as obtained in transmission SANS geometry in the crystalline phase. Several rings of increased intensity are visible. The intensity along the rings is constant, confirming a three dimesional powder structure. ˚ −1 , The circles in the map mark the position of the first five Debey-Scherrer rings ((111) at |Q| = 0.037 A −1 −1 −1 −1 ˚ ˚ ˚ ˚ (200) at |Q| = 0.043 A , (022) at |Q| = 0.060 A , (311) at |Q| = 0.071 A , (222) at |Q| = 0.074 A ) as ˚ and confirming earlier results from this sample [18]. expected for a fcc lattice with a lattice constant of 295 A However, due to the absence of sharp reflections relating to the powder structure in this scattering geometry it is not possible to decide whether the micelles indeed order in a fcc-like structure. Figure 18.19 shows the full transmision small angle scattering pattern for the sample with an applied shear rate of 100 s−1 . The shear gradient was perpendicular to the scattering plane along the z-direction, the flow was approximately along the x-direction and the y-direction was more or less along the vortex direction. In this representation it is visible

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Figure 18.18 Small angle scattering pattern for a fcc micellar powder structure. Reprinted from [17]. Copyright (2007) with permission from Elsevier.

Figure 18.19 Small angle scattering pattern for a fcc micellar crystal under shear. Reprinted from [17]. Copyright (2007) with permission from Elsevier.

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Figure 18.20 Grazing incidence small angle scattering pattern taken for the fcc micellar crystal. Reprinted from [17]. Copyright (2007) with permission from Elsevier.

that the (220) powder ring is condensed into intense Bragg peaks which exhibit a six-fold symmetry. The angular intensity distribution with six fold symmetry shows in addition that the [220] direction is oriented parallel to the flow velocity which was slightly tilted with respect to the x-direction [19]. As seen from the angular width of the reflections the crystal is not very well ordered and as already seen in the radial intensity distribution intensity from, e.g. the (111) and (311) reflections is projected into the detector plane [20]. The scattering pattern with six-fold symmetry is explained by a highly twinned fcc crystal structure, ABCABC and ACBACB stacking, as often found in colloidal or micellar systems [19]. Figure 18.20 shows data taken at rest and under grazing incident beam geometry. Without shearing the sample several well distinguishable peaks become visible for the same resolution as used for the transmission SANS measurement. The peaks are explained by a two dimensional powder structure with the [111] direction of the fcc lattice parallel to the surface normal (cubic dense packing). Note that in this case, next to the confining interface, really a two dimensional powder develops, which is in contrast to the twinned ABCABC and ACBACB layer sequence observed before and under shear. In the latter case Bragg reflections would only be visible if the crystal had by chance the proper orientation, i.e. (111) plane lies on the Ewald sphere. The Miller indices for a fcc lattice are shown in the map. Reflections indexed with a prime result from crystallites with an orientation that is tilted by 180◦ round the surface normal with respect to the unprimed indices. Now the (−111) and (200) reflections are, in addition to the small difference in |Q|, separated by an angle of 54.5◦ and even the (111) and (−111) reflections with equal |Q|-values become separated by 70.5◦ in angle. Indexing these spots is more straightforward and the structure can be resolved [17]. To obtain the real peak positions the Q-component parallel to the surface normal is corrected according to refractive effects. The effective

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 Q-values are given by the relation: Q = Q 2z,m − 4π (σl − σ Si ). Qz,m = 4π sin(θ ) denotes the measured Q λ value along the z direction as shown in Figure 18.6 and σ l and σ Si are the coherent scattering length densities of the liquid and the silicon, respectively.

18.4 Recent Experimental Results The field of neutron scattering for multiphase polymer systems is very wide and active. Accordingly, a complete review considering all aspects of the technique is very hard to provide and in the following the focus will be on some exemplary key experiments. These experiments offer new perspectives for the field and should provide further and detailed information on multiphase polymer systems. For the more conventional SANS studies the reader is referred to one of the reviews published recently [21–24]. In the following, first some spectroscopic experiments on polymer blends will be discussed. In the second part of the section the focus will be on diffraction from micelles. Micelles and block polymers are of great technical interest for the production of nano devices and structuring on mesoscopic length scales [2–6]. In particular, recent results will be exemplified on the three block copolymers called pluronics. These molecules have attracted considerable scientific interest [25–37]. In particular, the technique of grazing incidence small angle neutron scattering (GISANS) being an emerging technique [38] will be discussed in greater detail. 18.4.1

Polymer Dynamics

The investigation of dynamic miscibility in polymer blends, i.e., the question of how friction arises in chemically heterogeneous systems at present is a very active area of research [39–45]. In this context the question is how is a given dynamical process affected by blending and can we distinguish each blend component from a dynamical point of view, or do they move similarly when blended? It is noteworthy that, in addition to the technological importance of these systems, the investigation of the dynamical miscibility could also shed some light on the problem of the length scales associated with the different dynamical processes in polymers, and thereby contributing to the general understanding of the glass transition [46]. At temperatures well above the component glass-transition temperatures Tg the concept of segment self-concentration [39] provides a rather successful description of the component dynamics of a large number of polymer blends [42]. This self-concentration due to the chain connectivity is always enhanced locally and determines the component glasstransition behavior. On the other hand, in systems with greatly different component Tg ’s a decoupling of the dynamics of both components has been reported [poly(ethylene oxide)/poly(methyl methacrylate) (PEO/PMMA), polystyrene/poly(vinyl methyl ether), and PEO/poly(vinyl acetate)] [47–51]. In the system PEO/PMMA (TgPEO ≈ 200 K, TgPMMA ≈ 400 K), recent NMR studies have found up to 12 orders of magnitude different local relaxation times [48]. In such a situation, the low-Tg component moves in the random environment created by the frozen high-Tg component – a qualitatively different scenario compared to that where both components move on similar time scales [52]. Beyond its interest for blend dynamics including plasticizing effects, this situation may create an experimental test bed facilitating investigations of polymer chains in random environments, which presently was mainly studied theoretically [53, 54]. By use of neutron scattering on a miscible blend of two polymers with greatly different glass-transition temperatures Tg it has been shown that the nearly frozen high-Tg component imposes a random environment on the mobile chain. The results demand the consideration of a distribution of heterogeneous mobilities in the material and demonstrate that the larger scale dynamics of the fast component is not determined by the average local environment alone. This

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distribution of mobilities can be mapped quantitatively on the spectrum of local relaxation rates measured at high momentum transfers [55]. In another study, the question of the dynamic miscibility in the blend system composed by a mixture of 50% HHPP/50% PEP was studied. Exploiting isotopic labeling, the HHPP-component dynamics in the blend by using a mixture of protonated HHPP and deuterated PEP could experimentally be isolated. In the glassy state the methyl group dynamics of HHPP is not affected by blending. On the other hand, well above the average glass transition of the blend, the hydrogen motions in HHPP become faster in the presence of PEP. The combined analysis of these results and measurements on the fully protonated blend allows to deduce also the PEP-component dynamics in the blend in the supercooled liquid state. This dynamics is slowed down by blending but remains faster than that shown by the HHPP component in the blend [46].

18.4.2

Contrast Variation

Using contrast variation techniques in SANS measurements, the internal structure of micelles has been determined [56]. By this technique the micellar shape just above the critical micellization temperature or concentration was extracted. In particular, by using a range of D2 O to H2 O ratios in the solvent to highlight different regions of micelles formed by a three block copolymer (PEO-PPO-PEO), e.g. the PPO core, containing by far less water than the shell, and PEO corona regions, respectively, can be highlighted independently. This approach provides more direct information on the size and structure of the micelle core. Goldmints et al. have shown that the size of the PPO core of micelles formed by a PEO-PPO-PEO three block copolymer in aqueous solution is obtained directly from the form factor, which is very sensitive to the core dimension. In this case the assumption of equality of the corona and the hard-sphere interaction radii is not necessary, the model can be directly verified for the lower concentrated solutions, and the important micelle characteristics such as core and corona radii and water contents, and micelle aggregation numbers can be obtained directly for dilute solutions with no interactions between the micelles [56]. Another excellent and more unconventional example for the use of contrast variation has recently been presented by Lund et al. [57, 58]. In this experiment architectural and topological effects in the equilibrium chain exchange kinetics of polymeric micelles have been unraveled by the use of small angle neutron scattering. A study of micellar structures formed by poly(styrene)-poly(butadiene) (PS-PB) diblock copolymers and PB-PS-PB triblock copolymers in different n-alkane solvents is presented. Particular emphasis is placed on the dynamic properties of these micelles under equilibrium which are studied using a novel time-resolved small angle neutron scattering technique. The results show that the structures of the micelles are very similar for both the diblock and triblock copolymers, which allows a direct comparison of the dynamic properties. A novel logarithmic relaxation is found for both the triblock and the diblock micelles which is not consistent with theoretical expectations. However, for the diblock micelles, the relaxation kinetics seem to approach the rate and the single exponential decay predicted by [59] when the micellar cores are strongly swollen with solvent. For the triblock micelles a logarithmic relaxation is found for all cases as an effect of additional topological knots present even in highly swollen micellar cores. This behavior is assigned to an increased coupling of chain motion within the dense confined core – an effect which seems to vanish in diblock micelles when the core is sufficiently swollen.

18.4.3

Effect of Shear

Polymers in general are known to exhibit strong shear thinning. This is related to an orientation and disentaglement of the molecules and has been proven by neutron scattering and other techniques. Here the effect of shear on polymer blends and surfactant systems will be presented.

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Polymer Blends

A recent review on the effect of shear flow on the structure of multi-component polymer blends and solutions was given by Han et al. [60]. Older but more extensive ones have been presented by Larson [61] and Tyrrell [62]. In this context, mostly the techniques of small angle light and neutron scattering, optical microscopy, and fluorescence microscopy are used to directly assess the influence of an externally applied shear field on the phase stability and morphology of model polymer blends and solutions. The polymeric fluids of interest vary from miscible blends and pseudo-binary solutions near a critical point of unmixing to thermodynamically unstable and completely immiscible blends undergoing spinodal decomposition and coarsening in the presence of simple shear flow. Han et al. in particular review the influence that critical concentration fluctuations, viscoelasticity, and rheological asymmetry have on the shear response of polymer blends and solutions individually, and discuss the practically important interplay of these three separate effects. Neutrons have contributed important information in this context for many years. Boue et al. have presented small angle neutron scattering experiments on a semi-dilute polymer solution of fully deuterated linear polystyrene in protonated di-octyl-phthalate at rest and in laminar shear flow under a constant shear gradient in the range γ τ˙c > 1, with τ c being the longest relaxation time of the system [63]. An anisotropy of the two dimensional scattering pattern is observed under shear at all temperatures, with an increase of the low q scattering in the direction parallel to the flow. The corresponding double winged shape of the pattern resembles the butterfly shapes found for the interchain scattering in deformed gels and rubbers. Similar butterfly scattering patterns were also observed for crosslinked polystyrene solved in 1,2-dichloroethane [64]. Hobbie et al. [65] have used SANS in order to measure the influence of shear flow on a low-molecularweight polymer blend near the critical point. By combining with light scattering of the equilibrium critical dynamics, their measurements reveal that the long-range critical fluctuations begin to break apart when the shear becomes comparable to the characteristic relaxation rate, defined by the equilibrium lifetime of the critical fluctuations. The effect is found directly related to the decrease in the critical temperature caused by shear flow. Similar results have been reported by Hammouda et al. [66] Saito et al. [67] have performed a light scattering and SANS study in order to investigate the structural development under oscillatory shear flow in a semidilute polymer solution of high molecular weight deuterated polystyrene in di-octyl-phthalate. They observed the butterfly patterns in the small angle light scattering Q region under the low-frequency oscillatory shear flow and in the SANS Q region under the high-frequency oscillatory shear flow. Moreover, in the scattering profiles along the flow direction, the scattering peak due to the interdomain interference appeared. As the frequency increased, the scattering peak moved to the higher Q region. These results could be explained by the following interpretation: the higher the frequency, the shorter the time available for the structure formation in a shearing cycle, and hence the smaller the structures become. The SANS profiles could be reproduced assuming a shear-induced structure formation with a phase transition from the single-phase state to the two-phase state with a sharp interface [67].

18.4.3.2

Block Polymers in Solution

In this section several structural investigations on micellar systems will be described. Most studies are based on SANS techniques. Because of the large variety of experiments done with SANS on multiphase systems the focus is on several exemplary experiments. F¨orster et al. [68] have shown that shear thinning and orientation of wormlike micelles is observed for τdis γ˙ 1, where τ dis is the disentanglement time. They provide a direct relation between bulk properties like shear rate and shear viscosity, and molecular properties investigated by SANS such as micellar thickness and orientational distribution of long wormlike structures [68].

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Schmidt et al. have shown a similar shear orientation for a micellar hexagonal triblock copolymer phase [69]. On a microscopical length scale probed by SANS, the shear flow resulted in an alignment of rodlike PEO-PPO-PEO triblock copolymer micelles along the flow direction. This behavior is similar to that of polymer melts and solutions of low molar mass surfactants [70–73]. Small angle light scattering, however, provides information on the evolution of texture under flow. Here, scattering along the flow direction was observed corresponding to a stripe texture with stripes perpendicular to the flow direction [73]. The degree of shear alignment depends on strain, and rheo-birefringence data showed that simple shear and large amplitude oscillatory shear give the same results. The behavior of the lyotropic hexagonal phase of the PEO-PPO-PEO triblock copolymer under shear is thus very similar to that of low molar mass surfactants and block copolymer melts [69]. In another experiment Zipfer et al. probed the orientation of a lamellar phase under shear load. The mesoscopic structure in the lamellar region of ternary poly(ethylene oxide)-poly(propylene oxide)-poly(ethylene oxide) block copolymers and PEO-PPO-PEO/water/butanol systems was investigated by SANS. A transition from parallel to perpendicular alignment of lamellae, with a corresponding decrease in the viscosity, was found with increasing shear rate at high block copolymer concentrations. Additionally, a shear-induced formation of multilamellar vesicles was observed when the polymer concentration was reduced [74]. The effect of multilamellar micelle formation was studied in more detail by Escalante et al. investigating the effect of shear on a lamellar phase composed of the zwitterionic surfactant Cx DMAO (that became charged by 20 mol % through a protonation reaction) and a mediumchain alcohol as cosurfactant [75]. The shear-induced changes were studied by means of FF-TEM, and rheological, conductivity, and SANS measurements, and these experiments showed that the application of shear always leads to a transformation of the lamellar to multi-lammelar vesicle phase. From the experiments, three characteristic times have been assigned to the transformation process. The shortest time (only seen in the conductivity measurements) can be associated with the orientation parallel to the wall of the originally present lamellae in the shear field. The second time is the onset of the transformation of lamellae to vesicles and corresponds to a decrease of conductivity in flow direction and a minimum in the apparent viscosity. The third time is the one required to complete the conversion of planar lamellae to multilamellar vesicles and corresponds to reaching constant values for conductivity and viscosity. The authors found a similar behavior for a large variety of different systems and over a large range of shear rates. For a given system the characteristic parameter for the transformation process is the deformation applied. This means that such vesicle systems can be produced in a controlled manner and systems of desired properties can be obtained by adjusting the composition of the starting system and choosing the corresponding shear conditions [75]. Mortensen et al. have investigated aqueous solutions of triblock copolymers of poly(ethylene oxide)poly(propylene oxide)-poly(ethylene oxide), PEO-PPO-PEO by SANS. These systems offer a variety of phase behaviors [76]. In particular, the huge scattering contrast between the polymer and the solvent D2 O, for aqueous solutions makes neutron scattering the method of choice to extract information on the variety of phases and transitions between the various states appearing as a consequence of the strong change in the hydrophobic behavior of mainly PPO [77–80], and to a lesser extent also that of PEO, upon changes in temperature [81], pressure [82] and polymer concentration [81]. It has been shown by small angle neutron scattering that at low temperatures and/or concentrations the individual copolymers exist in solution as individual unimers. Thermodynamically stable micelles are formed with increasing copolymer concentration and/or temperature. The copolymer suspension undergoes as a result a first-order crystallization transition when the micellar volume fraction crosses the critical value for hard-sphere interaction. The crystalline powder sample changes abruptly into a monodomain material upon the application of shear [76]. Investigations on such shear-aligned monodomain samples reveal the crystallographic symmetry of the unit cell.

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Additionally, as result of an increasing micellar size upon increasing temperature, the micelles themselves may undergo a sphere-to-rod transition at elevated temperatures. In a shear field these rod-like micelles form a macroscopic nematic phase for low copolymer concentrations, and a hexagonal solid phase for higher concentrations. At even higher concentrations, lamellar mesophases are observed: one lamellar type is governed by the hydrophobic nature of PPO; one lamellar type appears at the highest copolymer concentration as a result of true crystallization of the PEO blocks [76]. Apart from aqueous solutions more complex systems have been investigated. For example, there have been studies of the effect of salts on the PEO-PPO-PEO micellization, see, e.g. [83–85], the complex association of PEO-PPO-PEO, low-molecular-weight surfactants, e.g. [86], the phase behavior of PEO-PPOPEO copolymers in mixtures of water and oil [87, 88] as well as the stabilization of oil–water emulsions [89–91]. Also the reverse kind of architecture with poly(ethylene oxide) as the middle block, i.e. PPO-PEO-PPO, exists and offers great new possibilities. Such copolymers can form micellar networks, which may provide the basis of attractive new properties [92]. It was noted early and shown above that soft materials that had been ordered and aligned in external fields exhibited Bragg peaks that could not be indexed assuming a single homogeneous crystal structure [93–95]. Their origin was attributed to changes in the crystal structure induced by the application of the external field. Such structural defects comprise changes in layer stacking, phase coexistence, twinning, unitcell distortions, specific multigrain arrangements or multiple diffraction [19, 96–98]. Field-induced structural defects would seriously prohibit the use of such methods to generate ordered soft materials for applications where homogeneous order and orientation are critical, such as in the preparation of photonic crystals or of nanostructured magnetic or semiconducting materials. F¨orster et al. have shown that such unexpected quasi-forbidden Bragg peaks are characteristic for ordered homogeneous macroscopically oriented soft materials. Their origin is related to the softness of the interaction potential, which on the one hand allows spontaneous ordering by tolerating imperfections of the constituent structures, but on the other hand also tolerates imperfections that limit the coherence of the crystalline lattice. This is shown by a combination of scattering experiments (neutron and synchrotron X-ray diffraction), direct imaging (scanning electron microscopy, SEM) and model calculations. It explains the presence of quasi-forbidden Bragg peaks in seemingly unrelated materials such as lyotropic liquid-crystalline phases, mesoporous materials, colloidal dispersions, block copolymers, electrorheological fluids and photonic crystals, as these materials are either soft or have been prepared from soft-material precursors. It is shown that a consistent picture of line density, field-induced orientation and epitaxial relations can be developed for the first time by the consideration of quasi-forbidden Bragg peaks [99]. 18.4.4

Near Surface Crystallization of Micelles

In the following, the near surface crystallization of a 18.5% (in weight) solution of the Pluronic F127 ((ethylene oxide)99 (propylene oxide)65 – (ethylene oxide)99 ) in deuterated water is discussed to exemplify the technique of grazing incidence neutron scattering. The bulk properties of this material have been reported in the literature in great detail [24, 37]. The solid wall was a SiO2 -terminated silicon wafer and showed a contact angle of water of around 50◦ . Figure 18.21 shows a typical result for a reflectivity (main panel) and GISANS (bottom right) measurement with the sample in the crystalline phase [100, 101]. In both cases the intensity is plotted over Qin-plane , defined in the plane of the interface, and Qout-of-plane , defined along the normal to the interface. The white and black color symbolizes high and low intensity, respectively. The specular reflectivity is found along the vertical dashed grey line. The area of total external reflection, indicated by high intensity, is visible close to Qin-plane = 0 and Qout-of-plane < 0.15 nm−1 . The pronounced reflection visible at Qout-of-plane ≈ 0.4 nm−1 is the first order

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Figure 18.21 Reflectivity and GISANS scattering pattern taken for the F127 sample in the fcc phase. Reprinted from [101]. Copyright (2009) with permission from American Chemical Society.

Bragg reflection on the specular line corresponding to the (111) reflection assuming a fcc structure (cubic close packing) [18]. Note, Qin-plane probes different length scales for the reflectivity and GISANS data. The neutron coherence length parallel to the projection of the incident beam on the interface can be up to 100 μm whereas it is only on the order of 10–100 nm along the perpendicular direction. The use of Qin-plane in both cases is justified, if a two dimensional random orientation of the crystallites in the plane of the interface is assumed [18]. Figure 18.22 shows rocking curves of the first order specular reflection taken in-plane and plotted for different temperatures. The micelles of this sample are known to form a cubic dense packing close to a solid interface with a lattice constant of 29.5 nm in the temperature range investigated [18]. Clearly, two components have to be separated. This is expected for scattering treated in the Born approximation with a cut of length

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Figure 18.23 The circles and squares, left panel, represent the correlation length and the intensity as obtained for a Gaussian fit to rocking curves not including the narrow line, respectively. The circles, right panel, depict the change in texture. Reprinted from [101]. Copyright (2009) with permission from American Chemical Society.

[14]. The narrow component (specular reflection) is related to scattering from the mean potential, averaged along the in-plane direction and over the coherence volume of the neutron beam, along the interface normal. The broad one (diffuse scattering) reflects the structural correlations resulting from fluctuations parallel to the interface which are on a length scale in the μm range [14]. For the narrow component a constant intensity and resolution limited line width is extracted for all temperatures. This is related to an unchanged layering (mean scattering potential) of the micelles along the normal of the interface. The circles in Figure 18.23, left panel, depict the correlation length as extracted from the broad component in the rocking curve plotted versus temperature. Two regimes can be separated. Between 20 and 35 ◦ C an exponential decay (black line) is found. At higher temperatures (35–55 ◦ C) the correlation length is constant, 1.63 ± 0.02 μm or about 55 unit cells. The squares in Figure 18.23, left panel, represent the intensity scattered diffusely into the (111) reflection. After an initial increase the reflection becomes weaker. The maximum matches well with the point when the correlation length reaches the minimum constant value. The circles in Figure 18.23, right panel, represent the difference of intensities (I) extracted from the GISANS pattern of the primed and the unprimed reflections divided by the sum and are related to the texture in the sample. Initially for temperatures around 20 ◦ C all reflections have the same intensity. At around 35 ◦ C texture develops. This is exactly the point where the correlation length changes from exponentially decaying to a constant value. This point also coincides with the maximum of the diffusely scattered intensity. Regarding this experimental facts, Figure 18.24 is a model illustration of how the crystallization in the investigated type of soft crystal proceeds. Dark and bright colors represent strong and weaker correlations, T = 20 °C

T = 21 °C

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T = 55 °C

Figure 18.24 Model for the crystallization of micelles at the solid–liquid interface. The circles symbolize crystallites. Dark color shows good correlations and lighter colour shows bad correlations. Reprinted from [101]. Copyright (2009) with permission from American Chemical Society.

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respectively. First crystallites are formed at different positions in the sample. Then the crystallinity changes in a two step process. First (20–35 ◦ C), the orientation of the crystallites remains random while they grow to sizes larger than 5 μm. With further increasing temperature more and more micelles agglomerate, the micelle core becomes more compact and the agglomeration number increases [56, 102]. These increasing potential fluctuations along the interface result in increased diffuse intensity. On the other hand, the changing micellar shape introduces strain enforcing a rearrangement of micelles and reducing the correlation length parallel to the interface. Second (35–55 ◦ C), the crystallites interpenetrate and start to rearrange their orientation. In addition, big crystallites take up the smaller ones. This process is known as Oswald ripening and results in the formation of a highly textured crystalline structure. The decreasing intensity in the diffuse component reveals that by the same time the fluctuations parallel to the interface become reduced. This happens at constant correlation length and could be related to anisotropic correlations resulting from Oswald ripening. No change in the mean potential (layering) is found over the whole temperature range indicating a well layered structure [101].

18.4.4.1

Effect of the Interface

To evaluate the effect of different surface energy on the crystallization, two single crystalline polished silicon disks were investigated as solid interfaces. One was oxidized for 15 min in a 5:1 mixture of H2 SO4 and H2 O2 resulting in a hydrophilic SiO2 termination (contact angle of water 33◦ ). The second substrate was cleaned with the same mixture of acids and a deposition by exposure to gaseous HMDS (1,1,1,3,3,3-Hexamethyldisiazane) was made for 24 hours to achieve a hydrophobic termination (contact angle of water 75◦ ). Figure 18.25 shows the specular reflected intensity taken with the sample in the liquid phase at 295 K at the hydrophilic (left panel) and the hydrophobic (right panel) interface. Both data sets are compared to ˚ the intensity the Fresnel-reflectivity. The data are corrected for the small angle scattering. At Qz = 0.04 A, is increased for the hydrophilic interface, showing layers of adsorbed micelles. A similar effect has been reported earlier [103]. For the hydrophobic interface the reflectivity drops faster than RF and no peak is visible. This shows a huge interface roughness and no adsorption of micelles. The inset in the left panel ˚ −1 , for the hydrophilic shows the rocking curve, including the small angle scattering, taken at Qz = 0.04 A

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Figure 18.26 GISANS data taken at a hydrophilic (panels a, b) and a hydrophobic interface (panels c, d) in the liquid phase at 295 K (panels a, c) and in the crystalline phase at 298 K (panels b, d). Reprinted from [18]. Copyright (2004) with permission from American Physical Society.

interface. The narrow peak at ω = 0 is resolution limited and represents in-plane correlations over a length scale exceeding 60 μm. The broad component arises from the liquid structure factor of the bulk sample. The asymmetry of the intensity is particularly well visible for off-set angles |ω| = 0.5◦ originates from the reduced penetration depth Dp (depth for which the intensity has dropped to 1/e) for negative ω-values. The dashed = 10−8 according to Ref. [13], with ρ the density and σ the sum of the line shows Dp calculated for β = λρσ 4π absorption and the incoherent scattering cross sections. Figure 18.26 shows GISANS data taken at an incident angle of 0.30◦ , slightly bigger than the total reflection edge (0.25◦ ), resulting in a penetration depth of the neutrons in the liquid of 40 μm. The panels on the top (a-b) and bottom (c-d) show data for the hydrophilic and the hydrophobic interface, respectively. For the left panels the sample is in the liquid phase at a temperature of 295 K. A ring of increased intensity ˚ −1 is visible for both interfaces, corresponding to the mean distance between micelles with radius 4 * 10−2 A ˚ of 160 A. At the hydrophilic interface sharp Bragg reflections become visible when the sample is heated to 298 K. The dots in the left panel of Figure 18.27 show the reciprocal lattice for a cubic close packed (ccp) structure with the [111] direction perpendicular to the interface. When the perpendicular axes [1–10] and [11–2] are randomly oriented around the unique direction [111] the dots become smeared out along the solid and dashed circles. This structure corresponds to a two dimensional powder known, e.g. from highly oriented pyrolytic graphite. Note that in our scattering geometry |k| ≈ 100 ∗ |Q 111 | and thus the Ewald sphere cuts through these powder rings (Figure 18.27). The diffraction peaks visible in Figure 18.26 are shown as solid circles in Figure 18.27, while reciprocal lattice points not probed by the experiment are shown as dashed lines. The data for the hydrophobic interface at 298 K feature four diffraction rings. Their absolute positions

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Figure 18.27 Reciprocal lattice for a cubic (ccp) and a hexagonal (hcp) close packing. The dots mark the positions of reciprocal lattice points. For a two dimensional powder with the unique axis along Qz perpendicular to the interface these points are smeared out into circles. On the right side the positions of the Bragg reflections in z-direction is shown with respect to the Ewald sphere. Reprinted from [18]. Copyright (2004) with permission from American Physical Society.

are explained by a three dimensional polycrystalline cubic face centered structure with a lattice constant of ˚ [18]. A similar structure was found earlier in the bulk [104]. 295 A 18.4.4.2

Near Surface Relaxation of Micelles

Within this experiment similar hydrophilic and hydrophobic coated silicon interfaces as during the previous measurement have been used. The scattering geometry is shown in Figure 18.28. To allow the shear experiment in a controlled manner a commercially-available rheometer (Bohlin CSR-10) was modified to fit the requirements of a neutron reflectometry study and mounted on the sample stage of the V6 reflectometer (Helmholtz Center Berlin for Materials and Energy, Germany). Neutrons collimated to a rectangular cross section of 6 * 1 mm2 enter a single crystalline silicon block, 70 * 70 * 10 mm3 , on the narrow side (70 * 10) and get then reflected at the bottom of the liquid sample. The left panel (a) in Figure 18.29 shows the Rocking curve (intensity normalized to the incident beam intensity) of the first order specular reflection taken in-plane with the Pluronic F127 dissolved in D2 O, exemplary against the hydrophilic silicon wafer. From Figure 18.29(a), it is clearly visible that shear has a strong influence on the line shape of the (111) reflection. An initially (without shear) narrow line representing

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Figure 18.28 Sketch of the experimental set up with the rheometer mounted on the sample stage of V6. Neutrons enter a single crystalline block of silicon on the narrow side and get scattered from the bottom of the solid–liquid interface. Reprinted from [100]. Copyright (2008) with permission from American Chemical Society.

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Figure 18.29 The panel on the left side (a) depicts Rocking curves for the (111) reflection measured for different shear rates and the panel on the right side (b) shows the change in the Rocking curve with time after stopping the shear with the sample at the hydrophilic wafer. Reprinted from [100]. Copyright (2008) with permission from American Chemical Society.

a lateral correlation length of some μm becomes clearly broadened under shear as seen from the reduced peak intensity. This is explained by a reduction in the correlation length and seems to be in contradiction to the known shear alignment of micellar crystals [105]. This fact can be explained by assuming a global alignment of crystallites under shear but at the same time a less well defined lattice constant as the sample is not in equilibrium and relates well to the explanation of secondary Bragg reflections given recently [99]. The right panel (b) in Figure 18.29 depicts the intensity (normalized to the incident beam intensity) distribution of the same reflection after stopping shear. Clearly, two components have to be separated. It turns out that the specular component increases with time while the diffuse one remains more or less unchanged. This is explained by the fact that the micelles form a more and more layered structure close to the interface without changing the lateral correlations or crystallite size. This manifests in the sharp and narrow peak visible at Qin-plane = 0 while the diffuse background remains constant. To extract information on the relaxation of the long-range orientational correlations (or layering) after stopping shear at the different interfaces the integrated intensity of the (111) reflection can be defined as order parameter. The result is depicted in Figure 18.30 where the order parameter (intensity normalized to the peak intensity before shearing the sample) for the two different interfaces is plotted versus time. The micelles rearrange on a time scale of several hours. The open and closed symbols represent data for the hydrophilic and

Order parameter

1.0 0.8 Hydrophobic surface Hydrophilic surface

0.6 0.4 0.2 0.0 0

2

4 Time [h]

6

8

Figure 18.30 Intensity of the narrow component (specular intensity) plotted versus time after stopping the shear taken with the sample in contact with a hydrophilic and a hydrophobic substrate. Reprinted from [100]. Copyright (2008) with permission from American Chemical Society.

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the hydrophobic interface, respectively. It turns out that the relaxation of the polymer micelles is dependent on the properties of the solid wall. For the hydrophobic interface a relaxation time of 2 hours is found whereas for the hydrophilic one 5 hours are extracted. In addition, the extrapolation for infinite times shows that only 25% of the initial (before shearing) intensity is recovered for the hydrophilic interface whereas at the hydrophobic one the crystal relaxes completely [100]. The micellar corona is more hydrophilic than the core and the interaction of the micelles with a hydrophilic interface is stronger [18, 103, 105]. This implies a tighter binding of micelles at this interface and explains the slower relaxation. On the hydrophobic wall, on the opposite, one expects weaker wall–micelle interactions and thus faster relaxation. 18.4.4.3

Dynamics Under Shear

In contrast to the studies presented above a 33% (in weight) solution of the three block copolymer P85 (EO25 – PO40 – EO25 ) in deuterated water was investigated. The phase diagram with respect to polymer concentration and temperature has been established by SANS [106]. Different phases from unimer over micelles to cubic and lamellar structures have been reported. Two single crystalline polished silicon wafers have been prepared as solid interfaces. One surface was oxidized by treatment of 15 min in Caro’s acid (5:1 mixture of H2 SO4 and H2 O2 ), which is known to form SiO2 at the interface with a hydrophilic property. The second silicon wafer was treated for one minute with HF. The surface coverage with hydrogen shows for times, longer than the duration of the experiment, the desired hydrophobic character. Figure 18.31 shows the scattering geometry for the neutron spectroscopy measurements. In reflection geometry (left panel) the vector of momentum transfer Q is parallel to the shear gradient and in transmission geometry Q is parallel to the flow of the liquid. Thus it is possible to distinguish the influence of shear on the microscopic dynamics in these both directions. In addition the inelastic Doppler scattered neutrons can for the transmission geometry provide information on the macroscopic flow [107–109]. The shear device dedicated for diffraction studies as well as that for the spectroscopy is described elsewhere in Ref. [110]. From a time of flight measurement, at the instrument NEAT (Helmholtz Center Berlin for Materials and Energy, Germany), of the sample mounted in a standard sample container the structure factor was extracted ˚ −1 . From the literature it is known that the maximum of the structure factor for heavy water peaking at 1.8 A ˚ −1 [111, 112]. The line width (right panel) for a Lorentzian fit to the data shows a Q2 is located at 2 A dependence revealing a translational diffusion. By assuming a random jump diffusion from the slope of the

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ow Fl

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Figure 18.31 Scattering geometry for neutron spectroscopy. The incoming and outgoing beam defines the scattering plane. The shear device is placed perpendicular to it. For the reflection geometry (left panel) Q is parallel to the shear gradient whereas for the transmission geometry (right panel) it is parallel to the macroscopic flow. Reprinted from [114]. Copyright (2005) with permission from European Physics Journal.

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Figure 18.32 Line width for the faster mode with the sample in reflection (left panel) and transmission (right panel) geometry at rest (closed circles) and for a shear rate of 6000 s−1 (open circles). Reprinted from [114]. Copyright (2005) with permission from European Physics Journal.

line width the diffusion constant can be calculated and is found to be 7.27 * 10−3 cm2 s−1 . This is lower than the literature data for the diffusion constant of 2.94 * 10−5 cm2 s−1 of deuterated water found at a temperature of 318 K [113]. The reduced diffusion constant observed in our study may most probably result from the restricted mobility of the water molecules bound to the hydration shell of the polymer micelles. Figure 18.32 reproduces the line width taken at an triple axis instrument with the shear device for Q-values between 0.3 ˚ −1 and for a temperature of 313 K. The left panel displays data taken with the sample in reflection and 1.7 A geometry and the right panel data taken with the sample in transmission geometry at rest and with an applied shear rate of 6000 s−1 . The quasielastic line widths that have been extracted from the data points are plotted versus Q2 . In addition the Q2 fit to the time of flight data discussed previously is shown for comparison as solid line. Two aspects have to be noted. 1. The diffusion of the water molecules in the sample is isotropic, in the two directions that have been addressed. 2. The diffusion is not affected by the applied shear. The inelastic Doppler scattered intensity from the flowing liquid is not visible in this experiment. The slower dynamics of the polymer monomers has been investigated separately and at different solid surfaces. From a standard measurement a translational jump diffusion but slower compared to that of the solvent was extracted and the diffusion constant was found to be 0.5 ± 0.1 * 10−6 cm2 s−1 . For a 10% solution of similar pluronics P104, P123 and F128 a diffusion constant of 2 * 10−7 cm2 s−1 has been found from light scattering at a temperature of 298 K for the monomer [81]. Both values are close to each other and the diffusive mode is assumed to be related to the monomer translation. Figure 18.33 shows quasielastic spectra taken with a 0.6 mm thick sample (10% scattering) in the shear ˚ −1 . The instrumental resolution of the shear device setup was device at a temperature of 290 K and Q = 1.4 A determined to be 1.4 μeV. For the measurements under shear the velocity of the moving surface was 1.3 m/s. The left panels show data taken with the hydrophilic interface as wall material of the moving and the fixed disk and the right panels show data taken with the hydrophobic interface. The two upper panels correspond to data taken with the sample at rest. The elastic line for the upper panels in reflection geometry (first line) as well as in transmission geometry (second line) is quasielastically broadened. The two spectra on the left side show a similar Lorentzian line width of 2 = 11.5 ± 1 μeV (FWHM) and 2 = 11.6 ± 1 μeV. With Fick’s law D = /(Q 2 ) the diffusion constant D = 4.4 ± 0.5 × 10−7 cm2 s−1 can be calculated. For the sample in contact with the hydrophilic interface the line narrows under shear (2200 s−1 ) to 7.0 ± 0.7 μeV (D = 2.7 ± 0.4 × 10−7 cm2 s−1 ) in reflection geometry but remains unchanged 11 ± 0.5 μeV in transmission geometry.

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geometry without shear

Reflection geometry under shear

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0.06 0.04 0.02 0.00

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FWHM = 11.6 µeV

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FWHM = 11 µeV linear

0.05 0.04 0.03 0.02

-15 -10 -5 0 5 10 Energy transfer (µeV)

15

Figure 18.33 Quasielastic spectra taken with a sample of a 33% (in weight) solution of P85 in deuterated water at a hydrophilic (left panels) and a hydrophobic interface (right panels). The top panels show data taken with the sample at rest either in reflection or in transmission geometry and the bottom panels show data taken at a shear rate of 2200 s−1 . Reprinted from [114]. Copyright (2005) with permission from European Physics Journal.

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This shows that under shear the diffusion constant becomes reduced in the direction of the shear gradient and becomes anisotropic. To fit the data in transmission geometry under shear the scattering law including the inelastic part from Doppler scattered neutrons was calculated. Due to the large quasielastic line width and the statistics a fit of the inelastic Doppler scattered part of the spectrum and macroscopic velocity profile was not possible [109]. For the hydrophobic interface shear has no influence on the quasielastic line width of 10.7 ± 1 μeV in reflection geometry [114].

18.5 Conclusion In this review the use of neutrons for multiphase polymer systems has been evaluated. The field is very active and broad and an exhaustive review is hardly possible. Accordingly, the focus is on the presentation of neutron scattering, in general as an introduction to the technique giving an idea of what possibilities are offered. Nearly all possible instruments are used in the field of multiphase polymer sytems, however, with a focus on SANS. Further, a couple of exemplary experiments are elucitated providing an introduction to the field. Particular interest is drawn to recent challenging experiments opening new fields for neutrons. In this context the focus is on the near surface crystallization and shear effects on the molecular ordering. It is shown that neutrons provide unique information complementing other experimental techniques.

Acknowledgements The author acknowledges financial support of the DFG (MA801/12-2 and ZA161/18-2) within the priority program (SPP) 1164 and the BMBF (ADAM 03ZA7BOC). Additionally, the author wants to thank Nicole Voss for her help in finalizing the manuscript.

References 1. M. Li and C. K. Ober, Block copolymer patterns and templates, Materials Today 9, 30 (2006). 2. S. Ludwigs, A. B¨ooker, A. Voronov, N. Rehse, R. Magerle, and G. Krausch, Self-assembly of functional nanostructures from abc triblock copolymers, Nat Mater 2, 744 (2003). 3. G. Register, On the straight and narrow, Nature 424, 378 (2003). 4. S. O. Kim, H. H. Solak, M. P. Stoykovich, N. J. Ferrier, de J. J. Pablo, and P. F. Nealey, Epitaxial self-assembly of block copolymers on lithographically defined nanopatterned substrates, Nature 424, 411 (2003). 5. C. D. Rosa, C. Park, E. L. Thomas, and B. Lotz, Microdomain patterns from directional eutectic solidification and epitaxy, Nature 405, 433 (2000). 6. J. Y. Cheng, C. A. Ross, E. L. Thomas, and et al., Fabrication of nanostructures with longrange order using block copolymer lithography, Appl. Phys. Lett. 81, 3657 (2002). 7. J. J. Sakurai, Modern Quantum Mechanics (Addison Weseley Publishing Company, 1994). 8. V. F. Sears, Thermal neutron scattering length and cross sections for condensed matter research (Theoretical Physics Branch Chalk River Nuclear Laboratories, 1984). 9. R. Scherm, Ann. Phys. 7, 349 (1972). 10. L. Van Hove, Correlations in space and time and born approximation scattering in systems of interacting particles, Phys. Rev. 95, 249 (1954). 11. H. Zabel, in Noncontractual Methods in Diffusion (Conf. Proc. of the Metallurical Society of AMIE; Pittsburgh, 1983). 12. T. Springer, Springer Tracts in Modern Physics, vol. 64 (Springer (Heidelberg), 1972). 13. L. G. Parratt, Surface studies of solids by total reflection of x-rays, Phys. Rev. 95, 359 (1954).

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14. S. Sinha, E. Sirota, S. Garoff, and H. Stanley, X-ray and neutron scattering from rough surfaces, Phys. Rev. B 38, 2297 (1988). 15. URL http://www.ill.eu/others/YellowBookCDrom/index.htm. 16. M. Wolff, K. Zhernenkov, and H. Zabel, The reflectometer adam at the ill: Present status and future perspectives, Thin Solid Films 515, 5712 (2007). 17. M. Wolff, A. Magerl, and H. Zabel, Ns-sans for the investigation of micellar systems, Thin Solid Films 515, 5724 (2007). 18. M. Wolff, U. Scholz, R. Hock, A. Magerl, V. Leiner, and H. Zabel, Crystallization of micelles at chemically terminated interfaces, Phys. Rev. Lett. 92, 255501 (2004). 19. W. Loose and B. J. Ackerson, Model-calculations for the analysis of scattering data from layered structures, J Chem Phys 101, 7211 (1994). 20. O. Diat, G. Porte, and J.-F. Berret, Orientation and twins separation in a micellar cubic crystal under oscillating shear, Phys. Rev. B 54, 14869 (1996). 21. Y. Melnichenko and G. Wignall, Small-angle neutron scattering in materials science: Recent practical applications, J. Appl. Phys. 102, 021101 (2007). 22. G. Wignall, Small Angle Neutron and X-Ray Scattering (Springer, 2007). 23. G. D. Wignall and Y. B. Melnichenko, Recent applications of small-angle neutron scattering in strongly interacting soft condensed matter, Rep. Prog. Phys. 68, 1761 (2005). 24. K. Mortensen, Structural properties of self-assembled polymeric aggregates in aqueous solutions, Polymer. Adv. Tech. 12, 2 (2001). 25. W. Brown, I. C. Schillh, M. Almgren, S. Hvidt, and P. Bahadur, Micelle and gel formation in a poly(ethylene oxide) poly(propylene oxide) poly(ethylene oxide) triblock copolymer in water solution - dynamic and static light-scattering and oscillatory shear measurements, J. Phys. Chem. 95, 1850 (1991). 26. P. Alexandridis and T. A. Hatton, Poly(ethylene oxide)-poly(propylene oxide)-poly(ethylene oxide) block-copolymer surfactants in aqueous-solutions and at interfaces - thermodynamics, structure, dynamics, and modeling, Colloid Surface A 96, 1 (1995). 27. P. Alexandridis, J. F. Holzwarth, and T. A. Hatton, Micellization of poly(ethylene oxide)-poly(propylene oxide)poly(ethylene oxide) triblock copolymers in aqueous-solutions - thermodynamics of copolymer association, Macromolecules 27, 2414 (1994). 28. K. Mortensen and J. S. Pedersen, Structural study on the micelle formation of poly(ethylene oxide) poly(propylene oxide) poly(ethylene oxide) triblock copolymer in aqueous-solution, Macromolecules 26, 805 (1993). 29. O. Glatter, G. Scherf, K. Schillen, and W. Brown, Characterization of a poly(ethylene oxide) poly(propylene oxide) triblock copolymer (eo(27)-po39-eo(27)) in aqueous-solution, Macomolecules 27, 6046 (1994). 30. K. Schillen, W. Brown, and R. M. Johnsen, Micellar sphere-to-rod transition in an aqueous triblock copolymer system - a dynamic light-scattering study of translational and rotational diffusion, Macomolecules 27, 4825 (1994). 31. K. Mortensen and Y. Talmon, Cryo-tem and sans microstructural study of pluronic polymer solutions, Macromolecules 28, 8829 (1995). 32. P. Alexandritis, V. Athanassiou, and T. A. Hatton, Pluronic-p105 peo-ppo-peo blockcopolymer in aqueous urea solutions - micelle formation, structure, and microenvironment, Langmuir 11, 2442 (1995). 33. P. Alexandridis, T. Nivaggioli, and T. A. Hatton, Temperature effects on structural-properties of pluronic p104 and f108 peo-ppo-peo block-copolymer solutions, Langmuir 11, 1468 (1995). 34. Q. F. Meilleur L, Hardy A, Probing the structure of pluronic peo-ppo-peo block copolymer solutions with their apparent volume and heat capacity, Langmuir 12, 4697 (1996). 35. P. Alexandridis and J. F. Holzwarth, Differential scanning calorimetry investigation of the effect of salts on aqueous solution properties of an amphiphilic block copolymer (poloxamer), Langmuir 13, 6074 (1997). 36. W. Loh, Block copolymer micelles, Encyclopedia Surface Colloid Sci. pp. 802–13 (2002). 37. K. Mortensen, W. Batsberg, and S. ren Hvidt, Effects of peo-ppo diblock impurities on the cubic structure of aqueous peo-ppo-peo pluronics . . . , Macromolecules 41, 1720 (2008). 38. P. Mueller-Buschbaum, R. Cubit, and W. Petry, Nanostructured diblock copolymer films: A grazing incidence small-angle neutron scattering study, Langmuir 19, 7778 (2003). 39. T. P. Lodge and T. C. B. McLeish, Self-concentrations and effective glass transition temperatures in polymer blends, Macromolecules 33, 5278 (2000).

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40. E. Leroy, A. Alegria, and J. Colmenero, Segmental dynamics in miscible polymer blends: Modeling the combined effects of chain connectivity and concentration fluctuations, Macomolecules 36, 7280 (2003). 41. Y. Hirose, O. Urakawa, and K. Adachi, Dielectric study on the heterogeneous dynamics of miscible polyisoprene/poly(vinyl ethylene) blends: Estimation of the relevant length scales for the segmental relaxation dynamics, Macomolecules 36, 3699 (2003). 42. Y. He, T. R. Lutz, and M. D. Ediger, Segmental and terminal dynamics in miscible polymer mixtures: Tests of the lodge-mcleish model, J. Chem. Phys. 119, 9956 (2003). 43. R. Kant, S. K. Kumar, and R. H. Colby, What length scales control the dynamics of miscible polymer blends, Macomolecules 36, 10087 (2003). 44. R. H. Colby and J. E. G. Lipson, Modeling the segmental relaxation time distribution of miscible polymer blends: Polyisoprene/poly(vinylethylene), Macomolecules 38, 4919 (2005). 45. E. Krygier, G. Lin, J. Mendes, and et al., Segmental dynamics of head-to-head polypropylene and polyisobutylene in their blend and pure components, Macomolecules 38, 7721 (2005). 46. R. P. Aparicio, A. Arbe, J. Colmenero, and et al., Quasielastic neutron scattering study on the effect of blending on the dynamics of head-to-head poly(propylene) and poly(ethylenepropylene), Macomolecules 39, 1060 (2006). 47. C. Lorthioir, A. Alegria, and J. Colmenero, Out of equilibrium dynamics of poly(vinyl methyl ether) segments in miscible poly(styrene)-poly(vinyl methyl ether) blends, Phys. Rev. E 68, 031805 (2003). 48. T. R. Lutz, Y. He, M. D. Ediger, and et al., Rapid poly(ethylene oxide) segmental dynamics in blends with poly(methyl methacrylate), Macomolecules 36, 1724 (2003). 49. A.-C. Genix, A. Arbe, F. Alvarez, and et al., Dynamics of poly(ethylene oxide) in a blend with poly(methyl methacrylate): A quasielastic neutron scattering and molecular dynamics simulations study, Phys. Rev. E 72, 031808 (2005). 50. M. Tyagi, A. Arbe, J. Colmenero, and et al., Dynamic confinement effects in polymer blends. a quasielastic neutron scattering study of the dynamics of poly(ethylene oxide) in a blend with poly(vinyl acetate), Macomolecules 39, 3007 (2006). 51. V. G. Sakai, C. Chen, J. K. Maranas, and et al., Effect of blending with poly(ethylene oxide) on the dynamics of poly(methyl methacrylate): A quasi-elastic neutron scattering approach, Macomolecules 37, 9975 (2004). 52. S. Hoffmann, L. Willner, and D. Richter, Origin of dynamic heterogeneities in miscible polymer blends: A quasielastic neutron scattering study, Phys. Rev. Lett. 85, 772 (2000). 53. G. Migliorini, V. G. Rostiasvili, and T. A. Vilgis, Polymer chain in a quenched random medium: slow dynamics and ergodicity breaking, Euro. Phys. J. B 33, 61 (2003). 54. B. Chakrabarti, Statistics of Linear Polymers in Disordered Media (Elsevier, Amsterdam, 2005). 55. J. Niedzwiedz, A. Wischnewski, M. Monkenbusch, and et al., Polymer chain dynamics in a random environment: Heterogeneous mobilities, Phys. Rev. Lett. 98, 168301 (2007). 56. I. Goldmints, G. E. Yu, C. Booth, K. A. Smith, and T. A. Hatton, Structure of (deuterated peo) (ppo) (deuterated peo) block copolymer micelles as determined by small angle neutron scattering, Langmuir 15, 1651 (1999). 57. R. Lund, L. Willner, J. Stellbrink, P. Lindner, and D. Richter, Logarithmic chain-exchange kinetics of diblock copolymer micelles, Phys. Rev. Lett. 96, 4 (2006). 58. R. Lund, L. Willner, D. Richter, and H. Iatrou, Unraveling the equilibrium chain exchange kinetics of polymeric micelles using small-angle neutron . . . , Applied Crystallography 40, S327 (2007). 59. A. Halpering and S. Alexander, Polymeric micelles: their relaxation kinetics, Macomolecules 22, 2403 (1989). 60. C. Han, Y. Yao, R. Zhang, and E. Hobbie, Effect of shear flow on multi-component polymer mixtures, Polymer 47, 3271 (2006). 61. R. G. Larson, Flow-induced mixing, demixing and phase transitions in polymeric fluids, Rheol. Acta 91, 497 (1992). 62. M. Tirrell, Phase-behavior of flowing polymer mixtures, Fluid Phase Equilibria 30, 367 (1986). 63. F. Boue and P. Lindner, Semidilute polymer-solutions under shear, Europhys. Lett. 25, 421 (1994). 64. E. Mendes, P. Lindner, M. Buzier, F. Boue, and J. Bastide, Experimental-evidence for inhomogeneous swelling and deformation in statistical gels, Phys. Rev. Lett. 66, 1595 (1991). 65. E. K. Hobbie, D. W. Hair, A. I. Nakatani, and C. C. Han, Crossover to strong shear in a low-molecular-weight critical polymer blend, Phys. Rev. Lett. 69, 1951 (1992). 66. B. Hammouda, A. I. Nakatani, D. A. Waldow, and C. C. Han, Small-angle neutron-scattering from deuterated polystyrene in dioctyl phthalate solution under shear, Macomolecules 25, 2903 (1992).

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67. S. Saito, S. Koizumi, K. Matsuzaka, S. Suehiro, and T. Hashimoto, Light scattering and small-angle neutron scattering studies of structures in a semidilute polymer solution induced under oscillatory shear flow, Macomolecules 33, 2153 (2000). 68. S. F¨orster, M. Konrad, and P. Lindner, Shear thinning and orientational ordering of wormlike micelles, Phys. Rev. Lett. (2005). 69. G. Schmidt, W. Richtering, P. Lindner, and P. Alexandridis, Shear orientation of a hexagonal lyotropic triblock copolymer phase as probed by flow birefringence and small-angle light and neutron scattering, Macromolecules 31, 2293 (1998). 70. K. Almdal, F. S. Bates, and K. Mortensen, Order, disorder, and fluctuation effects in an asymmetric poly(ethylenepropylene)-poly(ethylethylene) diblock copolymer, J. Chem. Phys. 96, 9122 (1992). 71. H. H. Winter, D. B. Scott, W. Gronski, S. Okamoto, and T. Hashimoto, Ordering by flow near the disorder-order transition of a triblock copolymer styrene-isoprene-styrene, Macomolecules 26, 7236 (1993). 72. T. Tepe, M. F. Schulz, J. Zhao, M. Tirrell, F. S. Bates, K. Mortensen, and K. Almdal, Variable shear-induced orientation of a diblock copolymer hexagonal phase, Macomolecules 1995, 3008 (1995). 73. G. Schmidt, S. Mueller, P. Lindner, C. Schmidt, and W. Richtering, Shear orientation of lyotropic hexagonal phases, J. Phys. Chem. B 192, 507 (1998). 74. J. Zipfel, J. Berghausen, G. Schmidt, and P. Lindner, Shear induced structures in lamellar phases of amphiphilic block copolymers, Phys. Chem. Chem. Phys. 17, 3905 (1999). 75. J. Escalante, M. Gradzielski, H. Hoffmann, and K. Mortensen, Shear-induced transition of originally undisturbed lamellar phase to vesicle phase, Langmuir 16, 8653 (2000). 76. K. Mortensen, Structural studies of aqueous solutions of peo-ppo-peo triblock copolymers, their micellar aggregates and mesophases; a small-angle neutron scattering study, J. Phys.: Condens. Matter 8, A103 (1996). 77. Y. Sasanuma, Conformational-analysis of poly(propylene oxide) and its model-compound 1,2- dimethoxypropane, Macomolecules 28, 8629 (1995). 78. C. Guo, H. Z. Liu, and J. Y. Chen, A fourier transform infrared study of the phase transition in aqueous solutions of ethylene oxide-propylene oxide triblock copolymer, Colloid and Polymer Science 277, 376 (1999). 79. C. Guo, J. Wang, H. Z. Liu, and J. Y. Chen, Hydration and conformation of temperaturedependent micellization of peo-ppo-peo block copolymers in aqueous solutions by ft-raman, Langmuir 15, 2703 (1999). 80. Y. L. Su, J. Wang, and H. Z. Liu, Ftir spectroscopic study on effects of temperature and polymer composition on the structural properties of peo-ppo-peo block copolymer micelles, Langmuir 18, 5370 (2002). 81. G. Wanka, H. Hoffmann, and W. Ulbricht, Phase-diagrams and aggregation behavior of poly(oxyethylene)poly(oxypropylene)-poly(oxyethylene) triblock copolymers in aqueoussolutions, Macromolecules 27, 4145 (1994). 82. K. Mortensen, D. Schwan, and S. Janssen, Pressure-induced melting of micellar crystal, Phys. Rev. Lett. 71, 1728 (1993). 83. P. Bahadur, K. Pandyn, M. Almgren, P. Li, and P. Stilbs, Effect of inorganic salts on the micellar behavior of ethylene oxide-propylene oxide block-copolymers in aqueous-solution, Colloid Polym. Sci. 271, 657 (1993). 84. J. H. Ma, C. Guo, Y. L. Tang, J. Wang, L. Zheng, X. F. Liang, S. Chen, and H. Z. Liu, Salt-induced micellization of a triblock copolymer in aqueous solution: A h-1 nuclear magnetic resonance spectroscopy study, Langmuir 23, 3075 (2007). 85. K. A. Dill, T. M. Truskett, V. Vlachy, and B. Hribar-Lee, Modeling water, the hydrophobic effect, and ion solvation, Annu. Rev. Biophys. Biomol. Struct. 34, 173 (2005). 86. E. Hecht, K. Mortensen, M. Gradzielsky, and K. Hoffmann, Interaction of aba blockcopolymers with ionic surfactants – influence on micellization and gelation, J. Phys. Chem. 99, 4866 (1995). 87. P. Alexandridis, U. Olsson, and B. Lindman, Phase behavior of amphiphilic block copolymers in water-oil mixtures: The pluronic 25r4-water-p-xylene system, J. Chem. Phys. 100, 280 (1996). 88. J. H. Ma, Y. Wang, C. Guo, H. Z. Liu, Y. L. Tang, and P. Bahadur, Oil-induced aggregation of block copolymer in aqueous solution, J. Chem. Phys. B 111, 11140 (2007). 89. J. S. Lettow, T. M. Lancaster, C. J. Glinka, and Y. J. Ying, Small-angle neutron scattering and theoretical investigation of poly(ethylene oxide)-poly(propylene oxide)-poly(ethylene oxide) stabilized oil-in-water microemulsions, Langmuir 21, 5738 (2005).

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90. P. Schmidt-Winkel, C. J. Glinka, and G. D. Stucky, Microemulsion templates for mesoporous silica, Langmuir 16, 356 (2000). 91. P. Holmqvist, P. Alexandridis, and B. Lindman, Modification of the microstructure in poloxamer block copolymerwater-“oil” systems by varying the “oil” type, Macomolecules 30, 6788 (1997). 92. K. Mortensen, W. Brown, and E. Jorgensen, Phase-behavior of poly(propylene oxide) poly(ethylene oxide) poly(propylene oxide) triblock copolymer melt and aqueous-solutions, Macromolecules 27, 5654 (1994). 93. J. Buitenhuis and S. Forster, Block copolymer micelles: Viscoelasticity and interaction potential of soft spheres, J. Chem. Phys. 107, 262 (1997). 94. T. Liu, C. Burger, and B. Chu, Nanofabrication in polymer matrices, Progr. Polym. Sci. 28, 5 (2003). 95. B. J. Ackerson, J. B. Hayter, N. A. Clark, and L. Cotter, Neutron-scattering from charge stabilized suspensions undergoing shear, J. Chem. Phys. 84, 2344 (1986). 96. P. Panine, T. Narayanan, J. Vermant, and J. Mewis, Structure and rheology during shearinduced crystallization of a latex suspension, Phys. Rev. E 66, 022401 (2002). 97. A. V. Petukhov, I. P. Dolbnya, D. G. A. I. Aarts, G. J. Vroege, and H. N. Lekkerkerker, Bragg rods and multiple x-ray scattering in random-stacking colloidal crystals, Phys. Rev. Lett. 90, 028304 (2003). 98. S. Forster, A. Timmann, M. Konrad, C. Schellbach, A. Meyer, S. Funari, P. Mulvaney, and R. Knott, Scattering curves of ordered mesoscopic materials, J. Phys. Chem. B 109, 1347 (2005). 99. S. F¨orster, A. Timmann, C. Schellbach, A. Fr¨oomsdorf, A. Kornowski, H. Weller, S. V. Roth, and P. Lindner, Order causes secondary bragg peaks in soft materials, Nat Mater 6, 888 (2007). 100. M. Wolff, R. Steitz, P. Gutfreund, N. Voss, S. Gerth, M. Walz, A. Magerl, and H. Zabel, Shear induced relaxation of polymer micelles at the solid-liquid interface, Langmuir 24, 11331 (2008). 101. M. Wolff, A. Magerl, and H. Zabel, Crystallization of soft crystals, Langmuir 25, 64 (2009). 102. J. S. Pedersen and M. C. Gerstenberg, The structure of p85 pluronic block copolymer micelles determined by small-angle neutron scattering, Colloids and Surfaces A: Physicochem. Eng. Aspects 213, 175 (2003). 103. M. C. Gerstenberg, J. S. Pedersen, and G. S. Smith, Surface induced ordering of micelles at the solid-liquid interface, Phys. Rev. E 58, 8028 (1998). 104. C. Wu, T. Liu, B. Chu, and et al., Characterization of the peoppopeo triblock copolymer and its application as a separation medium in capillary electrophoresis, Macomolecules 30, 4574 (1997). 105. E. Eiser, F. Molino, G. Porte, and O. Diat, Nonhomogeneous textures and banded flow in a soft cubic phase under shear, Phys. Rev. E 61, 6759 (2000). 106. K. Mortensen, Structural studies of aqueous solutions of peo-ppo-peo triblock copolymers, their micellar aggregates and mesophases; a small-angle neutron scattering study, J Phys-Condens Mat 8, A103 (1996). 107. M. Wolff, A. Magerl, B. Frick, and H. Zabel, Understanding of lubrication with-neutrons, Mat.-wiss. u. Werkstofftech. 34, 568 (2003). 108. M. Wolff, A. Magerl, B. Frick, and H. Zabel, Quasielastic neutron scattering for the investigation of liquids under shear, Chem. Phys. 292, 283 (2003). 109. M. Wolff, A. Magerl, B. Frick, and H. Zabel, Sheared liquids explored by means of neutron scattering, J. Phys.: Condens. Matter 15, S337 (2003). 110. M. Wolff, A. Magerl, R. Hock, B. Frick, and H. Zabel, Investigation of sheared liquids by neutron backscattering and reflectivity, Appl Phys A-Mater 74, S374 (2002). 111. J. A. Polo and P. A. Egelstaff, Neutron-diffraction study of low-temperature water, Phys. Rev. A 27, 1508 (1983). 112. F. Hajdu, S. Lengyel, and G. J. Palinkas, X-ray scattering and radial distribution function of liquid water, J. Appl. Cryst. 9, 134 (1976). 113. R. Mills, Self-diffusion in normal and heavy water in the range 1-45.deg., J. Phys. Chem. 77, 685 (1973). 114. M. Wolff, A. Magerl, and H. Zabel, Dynamics and structure in complex liquids under shear explored by neutron scattering, Phys. Rev. E 71, 13 (2005). 115. M. Wolff, A. Magerl, and H. Zabel, Structure of polymer micelles close to the solid interface, Euro. Phys. J. E 16, 141 (2005).

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19 Gas Diffusion in Multiphase Polymer Systems Eliane Espuche Ing´enierie des Mat´eriaux Polym`eres, UMR CNRS 5223, IMP@UCB, Universit´e de Lyon, Universit´e Lyon 1, France

19.1 Introduction The control of gas transport properties is a determining factor for a wide range of applications. Barrier properties are, indeed, required for packaging, protective coatings, fuel cell membranes whereas high permeability associated to high selectivity is needed for gas separation. Polymers are increasingly used in all these domains since they allow to achieve very different gas transport properties as a function of their structure and morphology. However, incremental improvements in polymer gas transport properties are still needed. In this context, multiphase polymer-based materials have become a major research and development topic as they could lead to a step-change in polymer performances. The aim of this chapter is to discuss the gas transport mechanisms in such systems and to draw the main factors that govern their properties. Two domains of interest have been considered: barrier multiphase materials and selective multiphase materials.

19.2 Gas Transport Mechanisms in Dense Polymer Films: Definition of the Transport Parameters Transport in dense polymer membranes is governed by a diffusion/solution process [1] which can be totally defined by three parameters. The diffusion coefficient (D) describes the kinetic aspect of the transport whereas the solubility coefficient (S) reflects the penetrant/polymer affinity and the thermodynamic aspect of the transport. The third parameter, the permeability coefficient (P), is a more global parameter which results from the contribution of both diffusion and solubility. Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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Figure 19.1 Gas transport through a dense polymer film of thickness l separating two compartments filled with a simple gas such as the gas pressure in the upstream compartment is p1 and the gas pressure in the downstream compartment is p2 .

Figure 19.1 describes the gas transport through a dense polymer film of thickness l separating two compartments filled with a simple gas such as the gas pressure in the upstream compartment is p1 and the gas pressure in the downstream compartment is p2 . The gas transport process can be divided into three steps: (i) adsorption of the molecules on the membrane upstream face (ii) diffusion/solution process of the molecules within the film (iii) desorption of the molecules from the downstream face. If we assume that adsorption and desorption equilibria are established instantaneously, the diffusion step becomes the key step of the transport process and the gas transport can be described according to the Fick’s first and second laws: J = −D

∂C ∂x

∂C ∂ 2C =D 2 ∂t ∂x

(19.1) (19.2)

where J is the gas flux, D is the diffusion coefficient and ∂C is the gas concentration gradient. ∂x In steady state, the gas flux J is constant and D can be deduced from J according to the following equation: J=D

(C1 − C2 ) l

(19.3)

where C1 and C2 are the gas concentrations in the polymer at the downstream and upstream faces of the membrane which are in equilibrium with the external pressures p1 and p2 , respectively.

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The equilibrium solubility coefficient of a gas in a polymer, S, is defined as the ratio of the concentration of gas dissolved in the polymer at equilibrium to the pressure of gas in the gas phase. S=

C p

(19.4)

The gas permeability represents the volumetric gas flow rate going, in steady state conditions, through a defined membrane area submitted to a fixed pressure gradient. The gas permeability coefficient is normalized on membrane thickness and it can be defined as: P=

Jl p1 − p2

(19.5)

Taking into account Eq. (19.3), the permeability coefficient can also be expressed as:  P=

C1 − C2 p1 − p2

 D

(19.6)

When the downstream pressure is much lower than the upstream pressure, Eq. (19.6) becomes: P=

C1 D p1

(19.7)

Thus in the case of a fickian transport, the three gas transport parameters are related by the following law: P = D.S

(19.8)

According to the free volume theory [2], diffusion in polymer media is supposed to occur through jumps from holes to holes of sufficient volume to accommodate the diffusing molecule. The concept of the cooperation of free volume is then particularly important in the understanding of diffusion in polymers. Indeed, the transport of small molecules depends on the segmental motion of polymer chains that can allow several holes to merge into one hole large enough for a diffusional jump to occur. From a structural point of view, the transport properties of a polymer will then depend (i) on the polymer chemical structure and on the intermolecular interchain interactions, (ii) on the polymer morphology, the crystalline phase being considered to be impermeable to small molecules, and (iii) on the polymer chain mobility. According to the diffusion/solution process, the affinity between the polymer and the diffusing molecule can also play a role in the transport phenomena, as it determines the solubility level. The solubility of a solute is as high as its critical temperature is important [3]. Solubility effects are then particularly important to be considered in transport of gases such as carbon dioxide and for water and organic vapours. High solubility can be at the origin of plasticization phenomenon (plasticization by water is, for example, commonly encountered with polar polymers, plasticization by CO2 can occur at high pressure) or even of anomalous transport behavior. All these considerations explain why very different transport properties can be achieved as a function of the polymer structure and morphology and as a function of the diffusing molecule. However, despite the large range of permeability coefficients that are covered by the different polymers, it is still difficult to meet all the requirements of a given application by using a lonely material. That is why multiphase systems are largely developed.

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Figure 19.2

Oxygen and water vapor permeability values for different polymers.

19.3 Multiphase Polymer Systems for Improved Barrier Properties 19.3.1

Introduction

Figure 19.2 reports the oxygen and water permeability values of different polymers that are commonly used in packaging applications. The data of Figure 19.2 clearly highlight that any polymer materials can be considered as a perfect barrier towards small molecules. Furthermore, except for polymers such as PVDC, polymers that are good gas barriers are poor water barriers and vice versa. It is then very difficult to find polymers with interesting barrier properties for polar and nonpolar diffusing molecules. This can be easily understood taking into account the transport mechanisms in polymer media. Polymers with high gas barrier properties have a polar structure that allows the formation of strong interchain interactions in the polymer amorphous phase. However, the polar groups that are distributed all along the polymer backbone chains have a high affinity towards polar molecules such as water, leading to an important water permeability. Furthermore, upon hydration, a plasticization phenomenon often takes place for these high hydrophilic materials and a loss of the initial polymer gas barrier properties is then observed [4]. Different ways have been considered either to improve initial polymer barrier properties or to keep high barrier levels upon hydration. They mainly consist of the association of materials with different properties leading to multiphase approaches. Three main types of morphology have been particularly studied. They consist of (i) dispersions of impermeable spheres within a polymer matrix, (ii) dispersions of impermeable lamellar nanofillers within a polymer matrix, and (iii) multilayer systems. For all these systems, the gas transport properties are governed by a pure diffusion/solution process and the transport mechanism is totally reversible. Another way has also been developed in order to keep, for a given period of time, a very high level of barrier properties towards a specific diffusing molecule. It consists in using active fillers that trap or react with the diffusing molecule. For these active membranes, the transport mechanism can no longer be considered as reversible. The gas transport mechanism in these different types of multiphase films will now be analyzed.

19.3.2

Dispersion of Impermeable Spheres Within a Polymer Matrix

One strategy that has been widely used to reinforce the barrier properties is to disperse an impermeable and non-miscible phase within the polymer matrix. The impermeable phase can either be composed of inorganic

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fillers or of a high barrier polymer. The most common morphology resulting from these blends consists of a dispersion of a nodular dispersed phase within a continuous matrix. Maxwell gave the earliest mathematical solution that allowed the description of the flow in such medium [5]. The assumptions of the model are:

r r r

the medium is considered to be a diluted medium so that there is no interaction between the dispersed particles the matrix local characteristics are not affected by the presence of the dispersed phase the polymer/filler interface is strong enough to avoid void creation at the interface.

Considering that the volume fraction of the dispersed phase is φd and the volume fraction of the continuous phase is φc , the general Maxwell law expression is: P − 1 = 3φd Pc



(Pd /Pc ) + 2 − φd (Pd /Pc ) − 1

−1 (19.9)

where P is the permeability of the blend, Pc the permeability of the continuous phase and Pd the permeability of the dispersed phase. If the dispersed phase is impermeable to small molecules, Pd = 0 and the Maxwell equation can be simplified. 2φc 2(1 − φd ) P = = Pc 3 − φc 2 + φd

(19.10)

A permeability law can also be established by considering the evolution of the solubility and the diffusion coefficients in the binary system with respect to the neat matrix. Indeed, the gas solubility in a medium composed of a dispersion of impermeable spheres within a polymer matrix can be expressed by: S = (1 − φd )S0

(19.11)

where S0 is the solubility coefficient in the neat polymer, S is the solubility coefficient in the composite, and φ d is the volume fraction of the impermeable dispersed phase. As far as the dispersed nodules act as impermeable obstacles, the diffusing molecules have to follow a more tortuous path to go through the composite film (Figure 19.3).

Figure 19.3 Representation of gas path through a composite film (d’) and gas path through the respective neat polymer (d).

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The tortuosity, τ , is defined as the ratio d /d and the diffusion rate is slow down in the composite material in comparison with the neat matrix. D can be expressed by Eq. (19.12): D=

D0 τ

(19.12)

where D0 is the diffusion coefficient in the pure polymer, D the diffusion coefficient in the composite and τ the tortuosity. If we assume that the transport mechanism is Fickian in the neat polymer and in the composite, it is possible to combine Eqs (19.8), (19.11) and (19.12) to give the expression of the relative permeability, i.e. the permeability of the composite ratioed to the permeability of the neat matrix: P (1 − φd ) = P0 τ

(19.13)

where P0 is the permeability coefficient in the pure polymer and P the permeability coefficient in the composite. The comparison of Eq. (19.10) with Eq. (19.13) leads to the following expression of the tortuosity, if we assume that the permeability of the continuous phase is equal to the permeability of the neat polymer. τ =1+

φd 2

(19.14)

According to this law, the tortuosity factor does not depend on the size of the dispersed spherical impermeable domains. It only depends on the total volume fraction of the impermeable dispersed phase. Furthermore the tortuosity does not differ as a function of the diffusing molecule. Bruggeman proposed another equation to calculate the permeability coefficient of a medium composed of a continuous phase containing spherical dispersed domains [5]. Assuming that the permeability of the continuous phase is Pc and that the permeability and volume fraction of the dispersed phase are Pd and φ d , respectively, Bruggeman defined the permeability of the blend, P, by the following equation:  (1 − φd ) =

Pd − P Pd − Pc



Pc P

1/3 (19.15)

Maxwell and Bruggeman equations lead to divergent results for high Pd /Pc ratio whereas they give quite similar results when the permeability of the dispersed phase is low in comparison with that of the continuous phase. For example, the relative permeability, P/Pc , is equal to 0.40 according to Maxwell law and to 0.35 according to the Brugggeman approach, for a blend in which 50% of the volume is composed of impermeable spheres (Pd = 0) [5]. These kinds of blends and morphologies have often been used to improve the barrier properties of low or medium gas barrier polymers, such as polyolefins or PET, respectively. Numerous works have been concerned with blending a polyolefin with polyamide (PA6 or PA66) or ethylene-vinyl alcohol copolymer (EVOH) [6–10]. Efforts to improve the gas barrier of PET by blending with EVOH [11], a liquid crystalline polyester [12] and aromatic polyamides [13, 14] have also been exploited. A third component, denoted as compatibilizer, is generally introduced at low content in these systems in order to avoid the development of weak interfaces between the two polymers and in order to improve the blend morphology and mechanical properties. Maxwell’s law has been shown to describe accurately the permeability of a great number of such systems [14–18]. In these systems, a volume fraction of impermeable dispersed phase equal to 0.2 generally leads to a decrease of the gas permeability of about 30%.

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Table 19.1 Expression of the relative permeability for different types of composites: composites formed by a uniform dispersion of impermeable spherical particles within a permeable matrix and composites composed of a dispersion of impermeable lamellar fillers placed perpendicular to the gas flux. P is the composite permeability, P0 is the permeability of the neat matrix, α is the filler aspect ratio. Homogeneous dispersion of impermeable spheres

Homogeneous dispersion of impermeable layers

Type of impermeable flakes

w t

Filler aspect ratio (α) P/ P0

19.3.3

1 1−φ φ 1+ 2 Maxwell law

α = w/t 1−φ αφ 1+ 2 Nielsen law

Influence of the Shape of the Dispersed Impermeable Phase: Interest of Oriented Polymer Blends and of the Nanocomposite Approach

According to Eq. (19.13), the increase observed in the barrier properties of a binary system strongly depends on the tortuosity. The tortuosity factor is directly related to the volume fraction and to the shape of the impermeable dispersed phase. Table 19.1 and Figure 19.4 allow the comparison of the dispersion of two types of impermeable fillers on the relative permeability: spheres and ribbons with different width-to-thickness ratios placed perpendicular to the gas flux.

Figure 19.4 Evolution of the relative permeability as a function of the filler volume fraction for fillers with different aspect ratios.

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Maxwell’s law has been used to describe the evolution of the relative permeability in the case of impermeable dispersed spheres whereas Nielsen’s law [19] has been developed for dispersions of impermeable ribbons lying in the plane of the film. In all cases, the relative permeability is expressed according to Eq. (19.13) and the tortuosity is defined by: τ =1+

αφ wφ =1+ 2t 2

(19.16)

where w is the filler width, t the filler thickness and α is the filler width-to-thickness ratio, namely the filler aspect ratio. For spheres the aspect ratio is equal to 1 and both the tortuosity and the relative permeability only depend on the filler volume fraction. The predicted increase of the gas pathway and of the tortuosity is quite low and Figure 19.4 represents the evolution of the relative permeability as a function of increasing amounts of impermeable spheres. The gas path increase can be much higher for a same volume fraction of impermeable ribbons. The relative permeability has been represented in Figure 19.4 for different values of α. For α equal to 2, the theoretical law becomes: 1 − φd P = P0 1 + φd

(19.17)

Equation (19.17) has been widely used to describe the properties of oriented semicrystalline materials [16, 20] or oriented polymer blends [21]. For these materials either the crystalline phase or the polymer that composes the dispersed phase are considered to be impermeable to small molecules. It is supposed that this impermeable phase can be represented by cylinders with a width-to-thickness ratio equal to 2 lying in the plane of the film. In this approach, it is also considered that the properties of the continuous phase are not affected by drawing. The curve representative of Eq. (19.17) underlines the interest of using oriented multiphase materials in comparison with isotropic multiphase materials for improved barrier properties (Figure 19.4). Numerous works concerning blends such as polyolefin/PA blends, polyolefin/EVOH blends or PET/PA blends have confirmed that higher barrier properties could be obtained by blend drawing [6–10, 14]. The high efficiency in reducing gas permeability is due to the fact that the high barrier constituent is dispersed as platelets oriented parallel to the direction of the gas flux. Furthermore, it has been shown that the aspect ratio of the dispersed platelets can vary in a large range depending on the biaxial draw ratio used to prepare the samples. It can, in particular, be much higher than the value of 2 considered in Eq. (19.17). For example, aspect ratio values as high as 27 have been obtained on a PET/aromatic PA blend (90/10 wt% composition). For this system, the relative permeability value has been shown to be equal to 0.36 [14]. According to this approach, it seems interesting to work with impermeable dispersed platelets of very high aspect ratios. Figure 19.4 shows that particularly interesting relative permeability values are calculated for lamellar platelets with a width-to-thickness ratio of a few hundreds. This platelet type is representative of inorganic lamellar nanofillers such as clays that have a thickness in a range of 1 nm and a width of few hundred nanometers. For these fillers, the theoretical evolution of the relative permeability as a function of the filler content can be divided in two steps: a high relative permeability decrease is observed in Figure 19.4 for low filler amount and then the relative permeability variation is much lower. Thus an optimum theoretical lamellar nanofiller content can be deduced for barrier properties. It is around 5% volume fraction. The theoretical decrease of the relative permeability as a function of the filler aspect ratio has been experimentally confirmed by Yano et al. [22] on nanocomposites based on polyimide and lamellar fillers with

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increasing width. In this study, all the fillers have the same thickness (1 nm) and the filler volume fraction is fixed to 2 wt%. The experimental results clearly show that the permeability decrease is higher as the filler width increases. Thus the tortuosity factor is, in this purely geometrical approach, the main factor that governs the barrier properties. It is then of great interest to understand how this factor can be affected by the filler dispersion state.

19.3.4 19.3.4.1

Nanocomposites Based on Lamellar Nanofillers Influence of the Filler Size and Orientation

Nielsen’s law has often been considered to calculate the theoretical permeability decrease in nanocomposite films with respect to the neat matrix. A good agreement has been observed between the theoretical and experimental relative permeability values for some systems [22, 23]. However, in most of the studies [24–32], the expected theoretical relative permeability values are lower than the experimental ones. The discrepancy between theoretical and experimental values is often attributed to a difference between the real morphology of the nanocomposite system and the ideal morphology considered in the Nielsen model (a regular array of ribbons of infinite length, of finite width w and thickness t lying perpendicular to the gas flux). Indeed, in a great majority of studies, the experimental filler aspect ratio value that allows to fit the experimental data is lower than the theoretical one. Different dispersion parameters, such as the degree of lamellar nanofiller delamination and the filler orientation, have then been considered and analyzed. Their impact on the inverse of the tortuosity that is, according to Eq. (19.13), proportional to the relative permeability is presented hereafter. Influence of the Degree of Lamellar Nanofiller Delamination [33] The dispersion of lamellae stacks leads, in comparison with a dispersion of individual lamellae, to a reduction of the mean filler aspect ratio and then, according to Eqs (19.13) and (19.16), to a decrease of the barrier properties. Figure 19.5 represents, for a filler volume fraction equal to 0.05, the evolution of the inverse of the tortuosity as a function of the filler

Figure 19.5 Evolution of the inverse of the tortuosity as a function of the filler aggregate thickness for different filler widths. The filler volume fraction is fixed to 0.05.

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stacks thickness for different filler stacks widths. The curves show that the filler aggregation phenomenon is all the less critical as the filler stacks width is high. Influence of the Filler Orientation A nonoptimized filler orientation leads to a decrease of the efficient width of the fillers and reduces the tortuosity. Bharadwaj [33] proposed to add an order parameter, S, to Nielsen’s equation to take account of the filler orientation. S is defined according to Eq. (19.18), where θ is the angle between the filler and the plane of the film. S=

 1 3 cos2 θ − 1 2

(19.18)

Different discrete values of S can be calculated as a function of the filler orientation. S is equal to 1 when the clays are perfectly lying in the plane of the film. It is equal to 0 for a random orientation of the clays within the film and it is equal to –1/2 when the clays are supposed to be placed perpendicular to the faces of the film. The tortuosity parameter of Nielsen’s law has been modified by Bharadwaj to take account of these different possible clay orientations: αφ τ =1+ 3



1 S+ 2

 (19.19)

Figure 19.6 presents the evolution of the inverse of the tortuosity as a function of the filler orientation for a filler volume fraction fixed to 0.05 and different filler widths. All curves underline that the highest barrier properties are obtained when the fillers are lying perpendicular to the gas flow. However, a defect in lamellar filler orientation is as less critical as the filler width is important. The general conclusion drawn from Bharadwaj’s analyses is that the quality of the filler dispersion in terms of orientation or exfoliation is as less crucial for barrier properties as the filler width is important.

Figure 19.6 Evolution of the inverse of the tortuosity as a function of the filler orientation for different filler width, w. The filler thickness is considered to be equal to 1 nm and the filler volume fraction is fixed to 0.05.

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19.3.4.2

759

Influence of the Filler Shape and Array

According to Fredrickson and Bicerano [34], Nielsen’s model, which assumes a regular array of ribbons, is accurate in the dilute regime but it is considered as inadequate by Cussler and co-workers [35] when the fillers concentration is semi-dilute, i.e. when the filler’s concentration is low but the fillers overlap (φ  1 but αφ  1). Thus new modeling has been recently proposed in order to take different types of array of the dispersed impermeable objects and also different shapes of these objects into account [35, 36]. Only models based on the dispersion of monodisperse impermeable flakes oriented perpendicular to the gas flux are presented hereafter [19, 34, 35, 37–39]. Table 19.2 summarizes for each considered model, the geometry of the dispersed fillers, the type of dispersion that is taken into account and the expression of the tortuosity factor τ . Three geometries of impermeable fillers have been considered, i.e. ribbons, hexagonal flakes and disks, leading to different values of the aspect ratio, α, defined in Table 19.2. Regular to random arrays of flakes are also considered. The expressions of the tortuosity that are presented in Table 19.2 are all derived from the equations of the relative permeability that take into account a reduction of the solubility due to the presence of impermeable flakes.

Table 19.2 Description of the different models based on the dispersion of monodisperse impermeable flakes in a permeable matrix. Model

Filler shape

Dispersion type

Aspect ratio α

Nielsen [10]

Ribbons of infinite length with a width w and a thickness t Ribbons of infinite length with a width w and a thickness t Ribbons of infinite length with a width w and a thickness t

Regular array

w/t

Regular array

w/t

Consideration of two courses of flakes with alignment and misalignment occurring with equal probability Random array

w/t

Random misalignment of successive layers of hexagonal flakes Random dispersion of non overlapping disks Random array

w/t

Cussler [27]

Cussler [27]

Lape and Cussler [26] Cussler [30]

Ribbons of infinite length with a width w and a thickness t Regular array of perfect hexagonal flakes

Gusev and Lusti [28]

Disk with a diameter D and thickness t

Fredrickson and Bicerano [25]

Disk with a diameter D and thickness t

w/t

D/t

D/t

Inverse of the tortuosity factor (1/τ ) 1

  αφ 1+ 2

  μ(αφ)2 1+ 4(1 − φ) with μ = 1    μ α2 φ2 1 1+ 4 1−φ with μ = 1/2 1

  1

αφ 1+ 3

2

   μ α2 φ2 1+ 4 1−φ with μ = 2/27  

 αφ 0.71 1 exp 3.47 2

  1 + x + 0.1245x 2 1 4 2+x with x = π αφ/2 ln(α /2)

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Relative Permeability

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0,8 0,6 0,4 0,2 0 0

0,02

0,04

0,06

0,08

0,1

Filler volume fraction Nielsen

Gusev-Lusti

Cussler: regular array

Cussler: random array

Cussler:hexagonal flakes

Fredrickson-Bicerano

Figure 19.7 Comparison of the relative permeability calculated from different models for a filler aspect ratio equal to 200 and filler volume fractions lower than 0.1.

Figure 19.7 gives a comparison of the relative permeability evolution calculated from these different models considering a filler aspect ratio value, α, equal to 200 and filler volume fractions, φ, lower than 0.1. These conditions are typical of those that are theoretically encountered with nanocomposites. From a general point of view, all these models indicate a very high reduction of the relative permeability for low filler contents and, excepting Cussler’s model which was developed for hexagonal flakes, the differences between all these models remain small. Some of these models, in particular the Cussler and Gusev & Lusti models, have been confronted with experimental data [24, 40, 41]. It seems that the barrier properties are often overestimated by all these new models, as in the case of Nielsen’s law. For all these models, as for Nielsen’s model, all the fillers are expected to have exactly the same aspect ratio. However, the dispersion of clays is often more complex, leading to the co-existence in a same nanocomposite film of different types of structures from exfoliated ones to intercalated ones [42, 43]. Additional models have then been further proposed to consider polydisperse flakes [35, 44]. 19.3.4.3

Dispersion of Polydisperse Impermeable Flakes

Cussler-Lape’s model [35] is a model based on polydisperse ribbons. In Cussler-Lape’s model [35], the ribbon thickness, t, is supposed to be constant and a discrete distribution of the ribbon width, w, is considered. The relative permeability is given by the following equation: 1−φ P =⎡ ⎛ ⎞ ⎤2 P0  1 φ ⎣1 + ⎝  ⎠ n i wi2 ⎦ 3 t n i wi i

(19.20)

i

A slight modification of Cussler-Lape’s model has been recently proposed by Picard et al. [44] in order to consider a variation of both the ribbon width and thickness. Indeed, in nanocomposites based on lamellar

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fillers, a modification of the filler aspect ratio is often related to the presence of filler stacks and then to a modification of the filler thickness. P 1−φ =⎡ ⎛ ⎤2 ⎞ P0  1 φ ⎣1 + ⎝  ⎠ n i (wi /ti )2 ⎦ 3 n i (wi /ti ) i

(19.21)

i

This law allows an accurate description of the experimental permeability data measured on PA6-based nanocomposites films prepared for a wide range of clay amounts (from 0 to 18 wt% of clay). In their work, Picard et al. [44] highlight also the necessity to consider not only the filler inorganic part but also the filler organic surfactant to calculate the filler volume fraction φ when filler stacks are observed within the polymer matrix. 19.3.4.4

Limitations of the Geometrical Models: Influence of Physicochemical Parameters

In the geometrical approach, the nanocomposite is always considered as an ideal binary system. It is supposed that the filler/matrix interactions are high enough to avoid voids creation at the polymer filler interface but low enough to avoid the formation of an interphase with specific properties. This hypothesis is not verified for all systems. Indeed, an increase of permeability has already been observed for some nanocomposites systems whereas higher reductions of permeability than those predicted by the geometrical laws have also been reported for other nanocomposites systems. Some authors [19, 45–47] have then proposed to take account of an interphase to model the gas transport properties. Xu [48] defined a factor of polymer chain-segment immobility to take account of a modification of the polymer chain mobility due to confined geometry environment. Beall [45] developed a model based on four contributions in the case of a strong interface. These contributions were linked to the impermeable fillers, the interface, the constrained polymer and the free polymer, respectively. At last, Wach´e [46] defined a model based on a three-phase system (the impermeable fillers, the polymer matrix and the interphase). This model assumes that the volume representative of the interphase is negligible but that the molecule diffusion rate in the interphase (Vi ) can differ from that in the polymer matrix (Vp ). Figure 19.8 schematizes the system. This model allows to describe the properties of nanocomposites with a strong or, on the contrary, with a weak interface. The tortuosity factor takes into account the interphase properties and it is defined by: τ =1+φ

  α  VP −1 1+ 3 Vi

(19.22)

Figure 19.8 Schematic representation of a nanocomposite system in which the diffusion rate at the filler/matrix interface is different from the diffusion rate in the bulk matrix.

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A strong interface appears as a higher density medium in which the diffusing molecules should be slow down. Hence, the Vp to Vi ratio should be higher than 1. On the contrary, a weak interface would be characterized by a higher mobility and a Vp to Vi ratio less than 1. Wach´e’s model has been successfully applied to liquid permeation in polyethylene based nanocomposites [46]. Thus in these new approaches, a decrease or an increase of the chain mobility in the filler vicinity can explain a deviation of the permeability values from the simple geometrical laws. However, all these models predict, as the more simple geometrical models, that the permeability decrease does not depend on the nature of the diffusing molecule. This trend is verified for a great majority of nanocomposites, however some systems depart from this law. For example, Ryu [49] reported on polyethylene-based nanocomposites’ different relative permeability values as a function of the gas nature. The dependence of the relative permeability on the type of diffusing molecule has been related by some authors to specific interactions between the clays or clay/polymer interphase and the diffusing molecule. Equation (19.11) that defines the solubility in nanocomposites has then been slightly modified: S = S0 (1 − βφ)

(19.23)

where β is a factor that depends on the interactions between the diffusing molecule and the clay or the clay/polymer interphase. This parameter can be negative as in the system (toluene/ polyolefin-montmorillonite composite) for which it has been estimated to be equal to –6.7 from sorption experiments [46]. 19.3.4.5

Gas Transport in Nanocomposites

All the different models that have been proposed in the literature to express the tortuosity factor as a function of the shape, orientation, dispersion state and volume fraction of dispersed impermeable particles agree to conclude that the barrier properties must be highly improved when lamellar nanofillers are individualized, uniformly dispersed within the polymer and are lying in the plane of the film. Furthermore, the models that take into account the role of the filler/matrix interface clearly show that a strong interface is favorable for optimized barrier properties. Due to its interesting aspect ratio (α ≈ 200) and to its high cation exchange capacity (CEC ≈ 1 mequiv/g), montmorillonite has been one of the most commonly-used swelling lamellar nanofillers for barrier properties and much attention has been paid to prepare from this filler, composites whose morphology is characterized by a high degree of clay exfoliation. The Na+ cation of the pristine montmorillonite is often exchanged by an organic surfactant in order to limit the filler hydrophilycity on the one hand and to favor clay/polymer interactions on the other hand. Among the different techniques developed to prepare polymer/clay nanocomposites, melt intercalation and in situ intercalative polymerization are particularly studied. Melt intercalation is an environmentally friendly preparation technique because no solvent is required and it can be applied to a wide variety of thermoplastic matrix [47, 50–55]. In situ intercalative polymerization consisting of the intercalation of a monomer into the swelling clay mineral followed by in situ polymerisation, has been largely used with thermosets [56–58] and has also concerned some thermoplastic polymers [40, 60–62]. Melt blending and in situ intercalative polymerization have been in particular applied to the preparation of nylon 6/montmorillonite composites or PCL/montmorillonite composites and exfoliated structures have been quite easily obtained with these matrices [62–72]. Due to the high clay delamination state and to the favorable clay/matrix interactions obtained in these systems, interesting gas barrier properties have been achieved. However, the barrier properties are often higher for composites prepared by in situ polymerization than for composites prepared by melt blending [62]. It was also demonstrated that the choice of specific clay surfactants could allow the grafting of initiators and the growth of polymer chains [40, 60, 62, 73] from the clay surface. By this way, both exfoliation of the

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platelets and strong clay/polymer interactions were favored, leading to higher reductions of the permeability than those that were predicted by the geometrical laws. If exfoliated clay structures are generally easily obtained with polar polymers, even by melt blending, it is not so obvious with apolar matrices such as polyolefins. Indeed, due to the low compatibility between the apolar matrix and the hydrophilic filler, a high degree of clay delamination is rarely obtained by melt blending a polyethylene with pristine clay. An increase of the permeability is even often observed for this composite system with respect to the neat matrix. The strategy to improve the clay dispersion state and the barrier properties is, first, to use an organomodified clay [24, 74–76] and, second, to introduce in the nanocomposite formulation modified polyolefins bearing polar groups, namely compatibilizers or interfacial agents [24, 26, 46, 74, 77–81]. Maleic anhydride grafted polyolefins have been one of the most widely-used interfacial agents [24, 26, 27, 46, 74, 77–81] for polyolefin-based nanocomposites. Generally, the increase of compatibilizer amount leads to an improvement of the clay dispersion state and to enhanced barrier properties [24]. Recently the compatibilizer family suitable for polyolefin has been enlarged to acrylic acid grafted polyolefins [46, 82], different copolymers [46, 83] or oxidized polyolefins [26, 79, 84]. Durmus et al. [79] and Picard et al. [84] studied the role of the interfacial agent polarity on the gas barrier properties. Both authors concluded that enhanced barrier properties were achieved when the concentration of polar groups in the interfacial agent increased. Picard et al. [84] concluded from their work that the compatibilizer polarity defined the strength of the clay/polymer interactions and the barrier properties. However, the relative permeability values reported in the literature for polyolefin-based nanocomposites [24, 26, 27, 82, 83, 84] are often limited (generally around 0.7) in comparison with the theoretical permeability decrease predicted by Nielsen’s law (around a factor 6 for 5% volume fraction of inorganic phase and a filler aspect ratio equal to 200). 19.3.5

Multilayers

Table 19.3 presents the gas permeability coefficients measured at 0%RH on some polar and nonpolar polymers and the evolution of the gas permeability as a function of the relative humidity. The gas permeability of polar polymers is generally very low at anhydrous state but it highly increases with the relative humidity [4, 85]. The gas permeability of nonpolar polymers is generally higher than that of polar polymers. However, due to their low water sorption capacity, these polymers keep their initial properties even at high relative humidity [4]. The interest of the multilayer approach is, from a gas transport point of view, to lead to materials that keep high gas barrier properties whatever the relative humidity. The most common multilayers are composed of synthetic polymers. The strategy consists of using a polar polymer such as EVOH as the internal layer and to protect this polymer from water by using a nonpolar polymer, generally a polyolefin, as external layers. An interfacial layer, often composed by a maleic anhydride grafted polyolefin, is also used in order to improve the polar and nonpolar layers’ adhesion. These assemblies are largely used in packaging applications

Table 19.3 Evolution of gas permeability coefficients as a function of the relative humidity for some polymers. Polymer

T(◦ C)

Gas

P at 0%RH (barrer)

Permeability variation

Polyethylene [4]

15 15 30 20 23

CO2 O2 CO2 CO2 CO2

6.1 1.6 0.1 0.01 0.001

Less than 5% between 0 and 43%RH 20% between 0 and 95%RH × 3 between 0 and 95% × 170 between 0 and 98%RH × 104 between 0 and 94%RH

Polyamide 6 [4] Chitosane [76] Polyvinyl alcohol [4]

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and a very thin metal layer is now often deposited on the assembly when a particularly high barrier level is required. Multilayers based in part or in their totality on biodegradable polymers have also been evaluated for food packaging [86, 87]. In particular, plasticized starches have been associated with polyethylene or with biodegradable polyesters (polycaprolactone, polylactic acid, polyesteramide, polyhydroxyalkanoates. . .). Some studies have also been focused on the association between chitosan and natural wax [88]. Packaging is not the only application domain of multilayers. These assemblies are, for example, also studied for fuel tank applications [89–91]. The strategy consists here again to combine polar and nonpolar polymer layers in order to limit the diffusion of both polar and nonpolar liquids (such as toluene/ethanol mixtures) and to reduce the emission of volatile organic compounds in the atmosphere. Equation (19.24) has been proposed to calculate the gas permeability of a multilayer composed of i layers whose volume fractions are φi and whose permeability coefficients are Pi .  φi 1 = P Pi i

(19.24)

This simple law has been used to predict the gas permeability of PEHD/PA6 multilayer and a good agreement has been observed between the experimental results and the calculated values for CO2 , O2 , N2 and water vapor [92]. However this equation cannot describe liquid or liquid mixture permeation. In that case, additional parameters have to be considered: the interactions between the permeate and the polymer, the role of the interface between two polymer layers and an eventual nonlinearity of the concentration gradient within the material. Equation (19.25) has been proposed by Del Nobile et al. [93] to calculate the permeability coefficient of a liquid in a three-layer system.   PwM awu , awd =

1 l1 1 1 l2 l3 1  + + ltot Pw,1 ltot Pw,2 aw1−2 , aw2−3 ltot Pw,3

(19.25)

where PwM (awu , awd ) is the average permeability of the three layer system, Pw,i is the permeability of the layer i, li is the width of the layer i, aw1−2 and aw2−3 are the activities at the interfaces, awu is the activity at the upstream face and awd is the activity at the downstream face. awd is supposed to be equal to 0. This equation has been successfully used to calculate the water permeability in polyolefin/EVOH/polyolefin multilayer [93]. However, this equation cannot be extended to the case of liquid mixtures. In fact, works are still in progress to model such complex systems. 19.3.6

Active Films

From a gas transport point of view, active films design barrier films that contain reactive additives, also named scavengers. These films are able to retard the breakthrough of penetrants from one side of the membrane to the other (Figure 19.9). The additive consumes the penetrant but it is also lost in the reaction. When the reaction is efficient, breakthrough occurs only after essentially all the additive is consumed. However, when enough solute has diffused into the membrane to exhaust the scavenger, solute can then diffuse steadily across the membrane as it would across an initially nonreactive film, with a largely unchanged permeability. These films have been widely developed for food packaging [94, 95] A mathematical model has been proposed by Yang et al. [96] to describe the effect of reactive additives on the breakthrough process. This model considers that penetrant moves through the membrane by normal

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te

of

the

pe

rm ea

Permeate amount

tio

nc urv e

Gas Diffusion in Multiphase Polymer Systems

As ym pto

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Figure 19.9

tL

t

Different permeation curves leading to the same time lag value.

Fickian diffusion, but upon encounter with additive, penetrant is lost by a reaction that is first order in both penetrant and additive concentrations. An important characteristic of this model is the Thiele modulus that is defined by Eq. (19.26). It reflects the relative rapidity of reaction processes in the membrane compared to diffusion across the membrane.  =

k R cl 2 D

(19.26)

where kR is the reaction constant, c: the concentration of the reactive additives within the film, l the film thickness and D the diffusion coefficient. A high Thiele modulus means that when the membrane is initially exposed to the penetrant, the penetrant is rapidly consumed when it encounters additives, and very little permeates through the membrane before nearly all the reactive additive is eliminated from the membrane. A low Thiele modulus, on the other hand, indicates that penetrant moves efficiently through the membrane, with small probability of reaction with additive. In this case, considerable amounts of penetrant may break through before all the additive is consumed. Intuitively, time lag which is the time–axis intercept of the asymptote of the permeation curve, as presented in Figure 19.9, is expected to be related to the Thiele modulus. One might expect that a higher Thiele modulus should lead to a longer time lag. However, numerical simulations reveal that the time lag is unaffected by the Thiele modulus when the initial distribution of additive in the membrane is constant. Siegel and Cussler [97] have in particular demonstrated that the time lag is independent of the reactivity of the additive, thus of the Thiele modulus, by considering a membrane of thickness l, in which the penetrant reacts irreversibly with an additive A, according to a fixed stoichiometry P + υA −−−−→ product

(19.27)

It is supposed that the film separates a reservoir containing a constant concentration, p ∗ , of penetrant and a perfect sink. It is supposed that the additive is dispersed uniformly throughout the membrane and that the additive loading is a0 . If H is the partition coefficient of penetrant between membrane and reservoir and D

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the diffusion coefficient, the time lag t L can be expressed by: l 2 a0 l2 l2 + tL = = 2υDH p ∗ 6D 6D



3a0 1+ υHp∗

 (19.28)

0 The time lag is predicted to increase by a factor of (1 + υ3a ) in comparison with the respective unreactive H p∗ film. However, this result is independent of the reaction rate of the scavenger. Siegel and Cussler [97, 98] have also considered the effect of a nonuniform dispersion of the additive within the film. The time lag expression becomes

tL =

1 υDHp∗



l

xa0 (x)d x +

0

l2 6D

(19.29)

When additive is dispersed uniformly throughout the membrane, a0 (x) = a0 , and the general Eq. (19.29) reduces to Eq. (19.28). More generally, Eq. (19.29) can be rewritten in the form: l2 tL = 6D

 1+

6 A0 x¯0 Hυp∗

 (19.30)

l l where A0 = 0 a0 (x)d x is the amount of initial additive per unit membrane area and x¯0 = 0 xa0 (x)d x/ l l 0 a0 (x)d x is the initial position of the centroidal plane with respect to additive mass. When additive is initially distributed uniformly, x¯0 = l/2. Since 0 < x¯0 < l, it follows that for a given amount of additive, the time lag is limited to values between zero and twice the value that would be measured when additive is uniformly distributed. The extremes correspond, respectively, to cases where all additive is crowded at the upstream or downstream face. According to this model, the time lag of reactive barrier membrane can be substantially increased by sequestering the scavenger closer to the downstream interface. This trend is supported by the experimental data obtained by Nuxoll et al. [98] on ZnO/PVA membranes. However, these authors show also that this improved time lag may sometimes be compromised by an increase in leakage across the membrane prior to time lag. The leakage increase becomes generally more dramatic as the scavenger is positioned very close to the downstream interface, where the solute concentration is lowest. 19.3.7

Comparison of the Different Ways Used to Improve Barrier Properties

Different ways can be used to improve barrier properties of polymers. They are all based on multiphase approaches. The most common way consists of the dispersion of an impermeable phase within the permeable continuous phase. The objective is then to obtain the morphology that leads to the highest tortuosity effect. In this way, the gain in barrier properties does not depend on the diffusing molecule and it has no limitation in time. The more efficient morphology consists in the dispersion of impermeable flakes with a high aspect ratio, perfectly individualized and lying in the plane of the film. A more specific approach can be developed when barrier properties are required for one specific molecule. In that case, reactive additives, also named scavengers, are dispersed within the polymer matrix. The resulting films are able to retard the breakthrough of the specific penetrant from one side of the membrane to the other. The additive consumes the penetrant but it is also lost in the reaction. The performances of these membranes are limited in time. When the reaction is efficient, breakthrough occurs only after essentially all additive is

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consumed. However, when enough solute has diffused into the membrane to exhaust the scavenger, solute can then diffuse steadily across the membrane as it would across an initially nonreactive film, with a largely unchanged permeability. The performances of the membrane highly depend on the concentration of reactive additives and on their location within the film. They are more efficient when they are sequestered close to the downstream face of the membrane. Finally, the multilayer approach is widely used when barrier properties are required in complex atmospheres. The association of polar and nonpolar polymer layers allows indeed to achieve good gas barrier properties at hydrated state or good barrier properties towards a mixture of polar and nonpolar liquids.

19.4 Multiphase Polymer-based Systems for Improved Selectivity 19.4.1

Introduction

Membranes are increasingly used for industrial separations and the separation of gases by using membranes is an attractive area due to potential energy savings in comparison to more traditional separation methods. Membrane transport properties are one of the determining parameters for this application: membrane permeability determines separation productivity and membrane selectivity determines separation efficiency. Membranes for gas separation are generally prepared by the solvent cast process. An ideal membrane is highly selective and very permeable. Unfortunately, polymers with a high selectivity are generally low permeable and vice versa. An upper-bound trade-off curve can be drawn considering permeability and selectivity of polymers for each type of gas separation [99]. It is observed that glassy polymers are generally located nearer to the upper-bound trade-off curve than rubbery polymers. Different ways have been investigated to advance the trade-off curve toward more selective, more permeable materials. The ideal selectivity for a couple of gases 1, 2 can be defined in a pure Fickian mechanism by: α21 =

P1 D1 S1 = P2 D2 S2

(19.31)

According to Eq. (19.31), the permeability selectivity results from a diffusion selectivity and a solubility selectivity. Due to the rather low gas solubility in polymers, the most common approach used to improve selectivity has been concerned with increasing diffusion selectivity. In particular, numerous studies have dealt with the design of glassy polymers characterized by free volumes of specific sizes. Polyimides have been one of the most widely-studied polymer families for this purpose [100–107]. Besides this approach, essentially concerned with polymer synthesis, a multiphase approach based on the association of organic and inorganic materials has also been developed. Some inorganic materials [108–111] display properties well above the trade-off curve for polymers. Dispersing such inorganic materials in polymers that already exhibit favorable performance on the trade-off curve could then provide a step-change in membrane properties. Thus organic-inorganic or ‘mixed matrix’ materials have become an important research topic in gas separation membrane science. Furthermore, these materials could combine the advantages of each medium: high separation capabilities of the inorganic fillers and the desired mechanical properties and economical processing capabilities of polymers. 19.4.2

Organic–inorganic Materials for Gas Separation Membranes

Zeolites [108, 109] and carbon molecular sieves [110, 111] offer very attractive permeation properties with permeabilities and selectivities significantly higher than polymer. They have been the most widely-used inorganic fillers for mixed matrix materials preparation.

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Zeolites are crystalline alumino-silicates composed of AlO4 and SiO4 tetrahedra, which build up a network of channels and cavities. The microcrystalline voids and channels which are interconnected are responsible for the very specific properties of these adsorbents. The aperture size is typically in the range of molecular ˚ Because the aluminium atom is trivalent, an excess of negative charge dimensions, between 3 and 10 A. is introduced in the network when Si is replaced by Al in the tetrahedral. This charge is compensated by non-framework cations, located near the negative charges. Because of the presence of cations, these zeolites are polar adsorbents. Therefore four main factors influence the properties of a zeolite:

r r r r

pore size which defines the ability of a molecule to enter and diffuse through the zeolite framework Si/Al ratio which determines the number of cations and thus the hydrophilicity of the zeolite framework type of cation (valence and size) direction of the pores (1, 2 or 3D porous network).

Carbon molecular sieves have also been used for gas separation applications. They are characterized by a pore size distribution which is much narrower than in the case of common active carbons. A definite pore size cannot be given for carbon molecular sieves, contrary to zeolites, but only a mean pore size which is also in the range of molecular dimensions. Carbon molecular sieves’ internal surfaces are basically hydrophobic, with possible variations due to acidic surface groups. The ideal morphology of organic–inorganic materials for high selective membranes should consist in the dispersion of inorganic fillers within the polymer matrix such as a continuous pathway through molecular sieves or zeolites should be obtained. This ideal morphology seems to be very difficult to obtain but it explains why high volume fractions of sieves are generally introduced in the polymer phase. Indeed, when discontinuities exist between the sieves, the gas is forced to diffuse through the polymer matrix that is a less performant medium. It is then very important to make an appropriate selection of the sieve and matrix phases. Having a highly permeable polymer matrix phase can result in bypass of the selective sieve phase, while a low-permeability polymer underutilizes the sieves and results in poor productivity. Increased performances have been generally obtained by the dispersion of high volume fractions (generally higher than 40%) of zeolites or carbon sieves in rubbery polymers [112, 113]. The results are not so successful with glassy polymers [114, 115]. Indeed, glassy polymer mixed matrix membranes often demonstrate poor polymer sieve adhesion [112, 116, 117], resulting in macroscopic voids and no selectivity enhancement. Different strategies have been proposed to improve polymer sieves’ adhesion: a modulation of the casting conditions and of the pre-and post-treatment applied to the membrane [115, 118], surface treatment of zeolites or use of chemically reactive techniques [119–121], use of plasticizers [118], and use of laboratorysynthesized carbon molecular sieve particles [122, 123]. Some interesting results have been obtained for some systems but they can not be generalized. In conclusion, the transport properties of organic–inorganic hybrid membranes are strongly dependent on the nanoscale morphology of the membranes. In particular, as for barrier multiphase systems, the characteristics of the polymer/filler interface represent a critical parameter for membrane performances. Moore and Koros [124] analyzed the effect of different types of interface on the resulting organic–inorganic membranes properties. According to these authors, Maxwell’s law (Eq. (19.32)) should allow the description of an ideal organic–inorganic membrane.  P = Pc

Pd + 2Pc − 2φd (Pc − Pd ) Pd + 2Pc + φd (Pc − Pd )

 (19.32)

where Pc is the permeability of the continuous phase (the polymer phase), Pd is the permeability of the dispersed phase (the zeolite phase) and φd is the volume fraction of the dispersed phase.

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A matrix rigidification around the sieves, generally underlined by an increase of the matrix glass transition temperature, should lead to a reduced permeability around the fillers. Moore et al. [124] proposed then to divide, in Maxwell’s law, the permeability of the continuous polymer phase Pc by a chain immobilization factor, β, to take account of this phenomenon. In that case, the permeability should decrease in comparison with the ideal mixed matrix without major changes in selectivity. In the case of a void at the interface, the permeability and the effective thickness of the void should be taken into account to modify Maxwell’s law. This morphology should lead to an increase of the permeability in comparison with the polymer matrix and, depending on the size of the void, to an unchanged or decreased selectivity in comparison with the neat polymer matrix. The last case should be met when the void thickness is in the order of the size of the gaseous penetrant. Stresses that arise at the organic–inorganic interface due to solvent evaporation during membrane formation should be one of the main causes that can explain this behavior. Permeabilities lower than those of the neat polymer with essentially no change in selectivity could be caused by the use of an impermeable zeolite or by some strongly-held sorbent that prevents the penetrant from entering the internal pores of the dispersed phase. The first case has been observed by Duval et al. [112]. These authors showed that the introduction of zeolite 5 A in a EPDM matrix led to a decrease of gas permeability without changes of selectivity. Zeolite 5 A behaved as impermeable silica dispersed within the matrix. The second case has been observed with some solvents (e.g. methanol or water) that can develop high interactions with hydrophilic zeolites and that can then greatly modify the access of small molecules to zeolite internal pores. In conclusion, the final properties of mixed membranes strongly depend on: (i) the initial properties of the organic and inorganic phases; (ii) the membrane preparation process; (iii) the membrane nanostructure; and (iv) the sieves/matrix interface.

19.5 Conclusion Multiphase polymer-based systems can allow significant improvement of gas transport properties in two main domains: barrier applications and gas selective membranes. However, to form successful multiphase materials, certain key requirements need to be met. First, particular attention has to be paid to the selection of the appropriate components to be combined. Simple laws can be used to estimate from the initial component properties, the expected properties of the multiphase system. The design of the multiphase material morphology is also a key factor for optimized properties. It must be adjusted as a function of the application domain. Thus, for barrier applications, homogeneous dispersion of impermeable lamellar nanofillers can lead to: a significant increase of initial barrier properties whatever the diffusing molecules; multilayers composed of polar and nonpolar polymers are essentially developed when barrier properties are required in complex atmospheres; and at last dispersion of active species with a preferred location near to the downstream face of the film is used when very high barrier properties towards a specific molecule are needed during a defined period of time. The ideal morphology for gas separation, on the other hand, consists of a homogeneous dispersion of molecular sieves or zeolites forming a continuous pathway through the sieves. Some laws have been proposed to model the properties of these different multiphase systems. These models have initially considered the multiphase material as an ideal binary system. However, the confrontations between experimental and theoretical data have often underlined significant differences. These discrepancies have been assigned either to a difference between the real and ideal morphology or to the formation of an interface with specific properties between the dispersed and the continuous phase. New modeling has then

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been developed to take account of these factors. In particular, it has been shown that the interface is a critical determinant of the overall transport properties and, from a practical point of view, most of the methodologies developed now to prepare multiphase polymer-based materials take care of the interface and tend to improve the interfacial properties by adapted film forming process, or by filler surface modifications.

References 1. J. Crank and G.S. Park, Methods of measurement, in Diffusion in Polymers, J. Crank, G.S. Park (Eds), Academic Press, New York, NY, 1968. 2. C.A. Kumins, T.K. Kwei, Free volume and other theories, in Diffusion in Polymers, J. Crank, G.S. Park (Eds), Academic Press, New York, NY, 1968. 3. R.T Chern, W.J. Koros, H.B. Hopfenberg, V.T. Stannett, in Material Science of Synthetic Membranes, D.R. Lloyd (Ed.), American Chemical Society, 1985. 4. J.A. Barrie, Water in polymers, in Diffusion in Polymers, J.Crank, G.S. Park (Eds), Academic Press, New York, NY, 1968. 5. R.M. Barrer, Diffusion and permeation in heterogeneous media, in Diffusion in Polymers, J.Crank, G.S. Park (Eds), Academic Press, New York, NY, 1968. 6. M.R. Kamal, I.A. Jinnah, L.A. Utracki, Permeability of oxygen and water vapour through polyethylene-polyamide films, Polym. Eng. Sci., 24, 1337–1347 (1984). 7. P.M. Subramanian, In Barrier Polymers and Structures, W.J. Koros (Ed.), Washington, DC: American Chemical Society; p. 252–265, 1990. 8. P.M. Subramanian, V. Mehra, Laminar morphology in polymer blends: Structure and properties, Polym. Eng. Sci., 27, 663–668 (1987). 9. S.Y. Lee, S.C. Kim, Laminar morphology development and oxygen permeability of LDPE/EVOH blends, Polym. Eng. Sci., 37, 463–475 (1997). 10. D. Jarus, A. Hiltner, E. Baer, Barrier properties of polypropylene/polyamide blends produced by microlayer extrusion, Polymer, 43, 2401–2408 (2002). 11. K.M. Kit, J.M. Schultz, R.M. Gohil, Morphology and barrier properties of oriented blends of poly-(ethylene terephtalate) and poly-(ethylene 2,6-naphtalate) with poly(ethylene-co-vinyl alcohol), Polym. Eng. Sci., 35, 680–692 (1995). 12. O. Motta, L. Di Maio, L. Incarnato, D. Acierno, Transport and mechanical properties of PET/Rodrun 3000 blown films, Polymer, 37, 2373–2377 (1996). 13. P.M. Subramanian, Poly(ethylene terephtalate) blends for permeability barrier applications, Polym. Eng. Sci., 27, 1574–1581 (1987). 14. Y.S. Hu, V. Prattipati, S. Mehta, D.A. Schiraldi, A. Hiltner, E. Baer, Improving gas barrier of PET by blending with aromatic polyamides, Polymer, 46, 2685–2698 (2005). 15. E. Espuche, M. Escoubes, J.P. Pascault, M. Taha, Transport properties of thermoplastic/thermoset blends, J. Polym. Sci. Part B: Polym. Phys., 37, 473–483 (1999). 16. L. Hardy, E. Espuche, G. Seytre, I. Stevenson, Gas transport properties of poly(ethylene 2,6 naphtalene dicarboxylate) films: Influence of crystallinity and orientation, J. Appl. Polym. Sci., 89, 1849–1857 (2003). 17. W.I. Higuchi, T. Higuchi, Theoretical analysis of diffusional movement through heterogeneous barriers, J. Amer. Pharm. Assoc., Scientific Edition, 49, 598–606 (1960). 18. C.M. Peterson, Permeability studies on heterogeneous polymer films, J. Appl. Polym. Sci., 12, 2649–2667 (1968). 19. L.E. Nielsen, Models for the permeability of filled polymer systems, J. Macromol. Sci. Chem. A 1, 929–942 (1967). 20. G.A.J. Orchard, I.M. Ward, The permeation of oxygen through drawn films of poly(ether ether ketone), Polymer, 33(19), 4207–4209 (1992). 21. R.M. Holsti-Miettinen, K.P. Perttila, J.V. Seppala, M.T. Heino, Oxygen barrier properties of polypropylene/polyamide 6 blends, J. Appl. Polym. Sci., 58, 1551–1560 (1995).

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22. K. Yano, K.A. Usuki, A. Okada, Synthesis and properties of polyimide-clay hybrid films, J. Polym. Sci. Part A: Polym. Chem., 35, 2493–2498 (1997). 23. D.J. Sekelik, E.V. Stepanov, S. Nazarenko, D. Schiraldi, A. Hiltner, E. Baer, Oxygen barrier properties of crystallized and talc filled poly(ethylene terephtalate), J. Polym. Sci. Part B: Polym. Phys., 37, 847–857 (1999). 24. J.E.P. Rupp, Polyethylene-Layered Silicate Nanocomposites, Z¨urich, Switzerland (2004). 25. T. Lan, P.D. Kaviratna, T.J. Pinnavaia, On the nature of polyimide-clay hybrid composites, Chem. Mater., 6, 573–575 (1994). 26. E. Jacquelot; E. Espuche, J.F. G´erard, J. Duchet, P. Mazabraud, Morphology and gas barrier properties of polyethylene-based nanocomposites P. J. Polym. Sci. Part B: Polym. Phys., 44, 431–440 (2006). 27. S. Hotta, D.R. Paul, Nanocomposites formed from linear low density polyethylene and organoclays, Polymer, 45, 7639–7654 (2004). 28. A.P. Shah, R.K. Gupta, H.V.S Gangarao, C.E. Powell, Moisture diffusion through vinyl ester nanocomposites made with montmorillonite clay, Polym. Eng. Sci., 42, 1852–1863 (2002). 29. C. Nah, H.J. Ryu, W.D. Kim, S.S. Choi, Barrier property of clay/acrylonitrile-butadiene copolymer nanocomposite, Polym. For Adv. Tech., 13, 649–652 (2002). 30. Y.P. Wu, Q.X. Jia, D.S. Yu, L.Q. Zhang., Structure and properties of nitrile rubber (NBR)-clay nanocomposites by co-coagulating NBR latex and clay aqueous suspension, J. Appl. Polym. Sci., 89, 3855–3858 (2003). 31. B. Brul´e, J.J. Flat, High barrier polyamide/polyolefin/organoclay nanocomposite, Macromol. Symp., 233, 210–216 (2006). 32. R.K. Bharadwaj, A.R. Mehrabi., C. Hamilton, C. Trujillo, M. Murga., R. Fan., A. Chavira, A.K. Thompson, Structure-property relationships in crosslinked polyester-clay nanocomposites, Polymer, 43, 3699–3705 (2002). 33. R.K. Bharadwaj, Modelling the barrier properties of polymer-layered silicate nanocomposites, Macromolecules, 34, 9189–9192 (2001). 34. G.H. Fredrickson and J. Bicerano, Barrier properties of oriented disk composites, J. Chem. Phys., 110, 2181–2188 (1999). 35. N.K. Lape, E.N. Nuxoll, E.L. Cussler, Polydisperse flakes in barrier films, J. Membr. Sci., 236, 29–37 (2004). 36. C. Yang, W.H. Smyrl, E.L. Cussler, Flake alignment in composite coatings, J. Membr. Sci., 231, 1–12 (2004). 37. A.A. Gusev, H.R. Lusti, Rational design of nanocomposites for barrier applications, Adv. Mater., 13(21), 1641–1643 (2001). 38. E.L. Cussler, S.E. Hughes, W.X.J. Ward, R. Aris, Barrier membranes, J. Membr. Sci., 38, 161–174 (1988). 39. G.D. Moggridge, N.K. Lape, C. Yang, E.L. Cussler, Barrier films using flakes and reactive additives, Progress in Org. Coat., 46, 231–240 (2003). 40. P.B. Messersmith, E.P. Giannelis, Synthesis and barrier properties of Poly(ε-caprolactone)-layered silicate nanocomposites, J. Polym. Sci., Part A: Polym. Chem., 33, 1047–1057 (1995). 41. S. Takahashi, H.A. Goldberg, C.A. Feeney, D.P. Karim., M. Farrell, K. O’Leary, D. R. Paul., Gas barrier properties of butyl rubber/vermiculite nanocomposite coatings, Polymer, 47, 3083–3093 (2006). 42. A. Vermogen, K. Masenelli-Varlot, R. S´egu´ela, J. Duchet-Rumeau, S. Boucard, P. Prele, Evaluation of the structure and dispersion in polymer-layered silicate nanocomposites, Macromolecules, 38, 9661–9669 (2005). 43. R. Xu, E. Manias, A.J. Snyder, J. Runt, New biomedical poly(urethane urea)-layered silicate nanocomposites, Macromolecules, 34, 337–339 (2001). 44. E. Picard, A. Vermogen, J.F. G´erard, E. Espuche, Barrier properties of nanocomposite membranes: influence of the clay content and dispersion state. Consequences on modelling, J. Membr. Sci., 292, 133–144 (2007). 45. G.W. Beall, New conceptual model for interpenetreting nanocomposites behaviour, Polymer-Clay Nanocomposites, 268–279 (2000). 46. R. Wach´e, Formulation et caract´erisation de poly´ethyl`enes charg´es avec des argiles-Propri´et´es barri`eres des nanocomposites obtenus, PhD, Universit´e de Bretagne Occidentale, France 2004. 47. L.M. Liu, Z.N. Qi, X.G. Zhu, Study on nylon 6 clay nanocomposites by melt intercalation process, J. Appl. Polym. Sci., 71, 133–1138 (1999). 48. R. Xu, Q. Zheng, Y. Song, Y. Shangguan, Calculating barrier properties of polymer/clay nanocomposites: Effects of clay layers, Polymer, 47, 2904–2910 (2006).

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49. S.H. Ryu, Y.W. Chang, Factors affecting the dispersion of montmorillonite in LLDPE nanocomposite, Polym. Bull., 55, 385–392 (2005). 50. R.R. Vaia, E.P. Giannelis, Lattice model of polymer melt intercalation in organically-modified layered silicates, Macromolecules, 30, 7990–7999 (1997). 51. M. Alexandre, G. Beyer, C. Henrist, R. Cloots, A. Rulmont, R. J´erˆome, P. Dubois, One pot preparation of polymer/clay nanocomposites starting from Na+ montmorillonite. 1. Melt intercalation of ethylene-vinyl acetate polymer, Chem. Mater., 13(11), 3830–3832 (2001). 52. P. Reichert, H. Nitz, S. Klinke, R. Brandsh, R. Thomann, R. M¨ulhaupt, Poly(propylene) organoclay nanocomposite formation: influence of compatibilizer functionality and organoclay modification, Macromol. Mater. Eng., 275, 8–17 (2000). 53. J.C. Huang, Z. Zhu, J. Yin, X. Qian, Y.Y. Sun, Poly(etherimide)/montmorillonite nanocomposites prepared by melt intercalation: morphology, solvent resistance and thermal properties, Polymer, 42, 873–877 (2001). 54. M. Laus, O. Francesangli, F. Sandrolini, New hybrid nanocomposites based on an organophilic clay and poly(styreneb-butadiene) copolymers, J. Mater. Res, 12, 3134–3139 (1997). 55. R.A. Vaia, S. Vasudevan, W Krawiec, L.G. Scanlon, E.P. Giannelis, New polymer electrolyte nanocomposites-melt intercalation of poly(ethylene oxide)in mica-type silicates, Adv. Mater., 7, 154–156 (1995). 56. Q.X. Zhang, K. Naito, B.Qi, Y. Kagawa, Epoxy nanocomposites based on high temperatue pyridinium-modified clays, J. Nanosci. Nanotechnol., 9, 209–215 (2009). 57. A. Di Gianni, E. Amerio, O. Monticelli, R. Bongiovanni, Preparation of polymer/clay mineral nanocomposites via dispersion of silylated montmorillonite in an UV curable epoxy matrix, Appl. Clay Sci., 42, 116–124 (2008). 58. V. Mittal, Epoxy-vermiculite nanocomposites as gas separation barrier, J. Comp. Mater., 42, 2829–2839 (2008). 59. K.M. Dean, S.A. Bateman, R. Simons, A comparative study of UV active silane-grafted and ion-exchanged organoclay for application in photocurable urethane acrylate nano- and micro-composites, Polymer, 48(8), 2231–2240 (2007). 60. L. Urbanczyk, C. Calberg, M. Alexandre, S. Benali, E. Espuche, F. Gouanve, S. Bourbigot, P. Dubois, C. J´erˆome, C. Detrembleur, Redispersion of PCL/clay masterbatches in situ synthesized in scCO2 into a poly(styrene-coacrylonitrile) matrix: morphology and properties, J. Mater. Chem., 18(39), 4623–4630 (2008). 61. J.H. Chang, Y.U. An, G.S. Sur, Poly(lactic acid) nanocomposites with various organoclays. I. Thermomechanical properties, morphology and gas permeability, J. Polym. Sci. Part B: Polym. Phys., 41, 94–103 (2003). 62. O. Gain, E. Espuche, E. Pollet, M. Alexandre, Ph. Dubois, Gas Barrier Properties of Poly(ε-caprolactone)/Clay Nanocomposites: Influence of the Morphology and Polymer/Clay Interactions, J. Polym. Sci. Part B, Polym. Phys., 43(2), 205–214 (2005). 63. S. Murase, A. Inoue, Y. Miyashita, N. Kimura, Y.J. Nishio, Structural characteristics and moisture sorption behavior of nylon 6/clay hybrid films, J. Polym. Sci. Part B : Polym. Phys., 40, 479–487 (2002). 64. Y. Kojima, A. Usuki, M. Kawasumi, A. Okada, T. Kurauchi, O.J. Kamigaito, Sorption of water in nylon-6 clay hybrid, J. Appl. Polym. Sci., 49, 1259–1264 (1993). 65. S.Y. Park, Y.H. Cho, R.A. Vaia; Three-dimensional structure of the zone-drawn film of the nylon-6/layered silicate nanocomposites, Macromolecules, 38, 1729–1735 (2005). 66. D.M. Lincoln, R.A. Vaia, Z.G. Wang, B.S. Hsiao, Temperature dependence of polymer crystalline morphology of nylon-6/layered silicate nanocomposites, Polymer, 42, 1621–1631 (2001). 67. T.D. Fornes, D.R. Paul, Crystallization behavior of nylon 6 nanocomposites, Polymer, 44, 3945–3961 (2003). 68. H.E. Miltner, G. Van Assche, A. Pozsgay, B. Pukanszky, B. Van Mele, Restricted chain segment mobility in polyamide 6/clay nanocomposites evidenced by quasi isothermal crystallization, Polymer, 47, 826–835 (2006). 69. P. Maiti, M. Okamoto, Crystallization controlled by silicate surfaces in nylon 6-clay nanocomposites, Macromol., Mat. Eng., 288, 440–445 (2003). 70. S. Xie, S. Zhang, H. Liu, G. Chen, M. Feng, H. Qin, F. Wang, M. Yang, Effects of processing history and annealing on polymorphic structure of nylon 6/montmorillonite nanocomposites, Polymer, 46, 5417–5427 (2005). 71. T.X. Liu, Z.H. Liu, K.X. Ma, L. Shen, K.Y. Zeng, C.B. He, Morphology, thermal and mechanical behavior of polyamide 6/layered silicate nanocomposites, Composites Sci. and Technol., 63, 331–337 (2003). 72. E. Devaux, S. Bourbigot, A. El Achari, Crystallization behavior of PA6 clay nanocomposite hybrid, J. Appl. Polym. Sci., 86, 2416–2423 (2002).

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73. B. Lepoittevin, N. Pantoustier, M. Alexandre, C. Calberg, R. J´erˆome, P. Dubois, Polyester layered silicate nanohybrids by controlled grafting polymerization, J. Mater Chem., 12, 3528–3532 (2002). 74. M. Kato, H.Okamoto, N.Hasegawa, A. Tsukigase, A. Usuki, Preparation and properties of polyethylene-clay hybrids, Polym. Eng. Sci., 43, 1312–1316 (2003). 75. M.A. Osmann., J.E.P. Rupp, U.W. Suter, Gas permeation properties of polyethylene-layered silicate nanocomposites, J. Mater. Chem., 15, 1298–1304 (2005). 76. E. Picard, H. Gauthier, J.F. G´erard, E. Espuche, Influence of the interlayer cations on the surface characteristics of montmorillonites. Consequences on the morphology and gas barrier properties of polyethylene/montmorillonites nanocomposites, J. Colloids Interf. Sci., 307, 364–376 (2007). 77. Y. Zhong, D. Janes, Y. Zheng, M. Hetzer, D. De Kee, Mechanical and oxygen barrier properties of organoclaypolyethylene nanocomposite films, Polym. Eng. Sci., 47, 1101–1107 (2007). 78. M.M. Reddy, R.K. Gupta, S.N. Bhattacharya, R. Parthasarathy, Structure-property relationship of melt intercalated maleated polyethylene nanocomposites, Korea-Australia Rheology Journal, 19, 133–139 (2007). 79. A. Durmus, M. Woo, A. Kasgoz, C.W. Macosko, M. Tsapatsis, Inytercalated linear low density polyethylene (LLDPE)/clay nanocomposites prepared with oxidized polyethylene as new type of compatibilizer: structural, mechanical and barrier properties, European Polymer Journal, 43, 3737–3749 (2007). 80. M. Mehrabzadeh, M.R. Kamal, Melt processing of PA-66/clay, HDPE/clay and HDPE/Pa-66/clay nanocomposites, Polym. Eng. Sc., 44, 1152–1161 (2004). 81. J.H. Lee, D. Jung, C.E. Hong, K.Y. Rhee, S.G. Advani, Properties of polyethylene- layered silicate nanocomposites prepared by melt intercalation with PP-g-MA compatibilizer, Composites Sci. and Technol., 65, 1996–2002 (2005). 82. S. Sanchez Valdes, M.L. Lopez Quintanilla, E. Ramirez Vargas, F.J. Medellin Rodriguez, J.M. Gutierrez Rodriguez, Effect of ionomeric compatibilizer on clay dispersion in polyethylene/clay nanocomposites, Macromol. Mater. Engineer., 291, 128–136 (2006). 83. R.K. Shah, R.K. Krishnaswamy, S. Takahashi, D.R. Paul, Blown films of nanocomposites prepared from low density polyethylene and a sodium ionomer of poly(ethylene-co- methacrylic acid), Polymer, 47; 6187–6201 (2006). 84. E. Picard, A. Vermogen, J.F. G´erard, E. Espuche, Influence of the compatibilizer polarity and molar mass on the morphology and the gas barrier properties of polyethylene/clay nanocomposites, J. Polym. Sci. Part B : Polym. Phys., 46(23), 2593–2604 (2008). 85. S. Despond, E. Espuche, A. Domard, Water sorption and permeation in chitosan films. Relation between gas permeability and relative humidity, J. Polym. Sci. Part B : Polym. Phys., 39, 3114–3127 (2001). 86. P. Dole, L. Averous, C. Joly, Evaluation of starch-Pe multilayers: Processing and properties, Polym. Eng. Sci., 45(2), 217–224 (2005). 87. L. Averous, Biodegradable multiphase systems based on plasticized starch: A review, J. Macromol. Sci.- Polym. Rev., 44(3), 231–274 (2004). 88. S. Despond, E. Espuche, N. Cartier, A. Domard, Barrier properties of paper-chitosan and paper-chitosan-carnauba wax films, J. Appl. Polym. Sci., 98(2), 704–710 (2005). 89. K. Kimura, T. Mashimo, M. Fukushima, N. Osawa, S. Komata, Multilayer plastic structures with high mechanical strength for prevention of permeation of nonpolar and polar materials, and their multilayer containers and fuel tanks, JP2005246673 (2005). 90. J.J. Flat, G. Bellet, B. Brul´e, P. Blondel, Multilayer structure having one layer based on polyamide and HDPE, EP1607213 (2005). 91. P. Delbarre, Fuel tank of multilayer structure having an exposed barrier layer, US200319876 (2003). 92. C.H. Huang, J.S. Wu, C.C. Huang, Predicting the permeability and tensile behaviour of high density polyethylene/tie/polyamide 6 three-layer films, Polym. Intern. (Tokyo Japan), 36, 386–393 (2004). 93. M.A. Del Nobile, G.G. Buonocore, D. Dainelli, L. Nicolais, A new approach to predict the water transport properties of multilayer films intended for food-packaging applications, J. Food Sci., 69 FEP85–FEP90 (2004). 94. K.F. Finger, A.P. Lemberger, T. Higuchi, L.W. Busse, D.E. Wurster, Investigation and development of protective ointments, J. Am. Pharm. Assoc., 49, 569–573 (1960). 95. L. Vermeiren, F. Devlieghere, M. Van Beest, N. De Kruif, J. Debevere, Developments in the active packaging of foods, Trends Food Sci. Technol., 10, 77–86 (1999).

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96. C. Yang, E. Nuxoll, E.L. Cussler, Reactive barrier films, AIChE J., 47, 295–302 (2001). 97. R.A. Siegel, E.L. Cussler, Reactive barrier membranes: some theoretical observations regarding the time lag and breakthrough curves, J. Membr. Sci., 229, 33–41 (2004). 98. E.E. Nuxoll, R.A. Siegel, E.L. Cussler, Layer reactive barrier films, J. Membr. Sci., 252, 29–36 (2005). 99. L. Robeson, Correlation of separation factors versus permeability for polymeric membranes, J. Membr. Sci., 62, 165–195 (1991). 100. Piroux F., Espuche E., Mercier R., Pineri M., G. Gebel Gas transport mechanism in sulfonated polyimides. Consequences on gas selectivity, Journal of Membr. Sci., 209 (1), 241–253 (2002). 101. M.R. Coleman, W.J. Koros, Isomeric polyimides based on fluorinated dianhydrides and diamines for gas separation applications, J. Membr. Sci., 50, 285–297 (1990). 102. S.A. Stern, Polymers for gas separations: the next decade, J. Membr. Sci., 94, 1–65 (1994). 103. M. Langsam, W.F. Burgoyne, Effects of diamine monomer structure on the gas permeability of polyimides. I. Bridged diamines, J. Polym. Sci., 31, 909–921 (1993). 104. K. Tanaka, Y. Osada, H. Kita, K.I. Okamoto, Gas permeability and permselectivity of polyimides with large aromatic rings, J. Polym. Sci.Part B: Polym. Phys., 33, 1907–1915 (1995). 105. M. Langsam, In: Polyimides, Fundamentals and Applications; Gosh, M.K., Mittal, K.L., Eds.; Marcel Dekker: New York, 1996; Chapter 22, p. 697. 106. M. Al-Masri, D. Fritsch, H.R. Kricheldorf, New polyimides for gas separation.2. Polyimides derived from substituted catechol bis(etherphthalic anhydrides)s, Macromolecules, 33(19), 7127–7135 (2000). 107. H. Jianhua Fang, H. Kita, K.I. Okamoto, Hyperbranched polyimides for gas separation applications. 1. Synthesis and characterization, Macromolecules, 33(13), 4639–4646 (2000). 108. J. K¨arger, D.M. Ruthven, Diffusion in Zeolites and Other Microporous Solids, WileyI Interscience, New York, 1992. 109. D.W. Breck, Zeolite Molecular Sieves: Structure, Chemistry and Use, John Wiley & Sons Inc., New York, NY, 1974. 110. H. Suda, K. Haraya, Carbon molecular sieve membranes: preparation, characterization, and gas permeation properties, In Proceedings of the ACS Symposium series 744 on the membrane formation and modification, American Chemical Society, 295–313 (2000). 111. D.Q. Vu, W.J. Koros, S.J. Miller, High pressure CO2 /CH4 separation using carbon molecular sieve hollow fiber membranes, Ind. Eng. Chem. Res., 41, 367–380 (2002). 112. J.M. Duval, B. Folkers, M.H.V. Mulder, G. Desgrandchamps, C.A. Smolders, adsorbent filled membranes for gas separation. Part I: Improvement of the gas separation properties of polymeric membranes by incorporation of microporous adsorbents, J. Membr. Sci., 80, 189–198 (1993). 113. M. Jia, K.V. Peinnemann, R.D. Behling, Molecular sieving effect of the zeolite-filled silicone rubber membranes in gas permeation, J. Membr. Sci., 57, 289–296 (1991). 114. T.M. G¨ur, Permselectiviy of zeolite filled polysulfone gas separation membranes, J. Membr. Sci., 93, 283–289 (1994). 115. M.G.M. S¨uer, N. Bac, L. Yilmaz, Gas permeation characteristics of polymer-zeolite mixed matrix membranes, J. Membr. Sci., 91, 77–86 (1994). 116. J.P. Boom, I.G.M. Punt, H. Zwijnenberg, R. de Boer, D. Bargeman, C.A. Smolders, H. Strathmann, Transport through zeolite filled polymeric membranes, J. Membr. Sci., 138, 237–258 (1998). 117. J.F. Vankelecom, E. Merckx, M. Luts, J.B. Uytterhoeven, Incorporation of zeolites in polyimide membranes, J. Phys. Chem., 99, 13187–13192 (1995). 118. R. Mahajan, R. Burns, M. Schaeffer, W.J. Koros, Challenges in forming successful mixed matrix membranes with rigid polymeric materials, J. Appl. Polym. Sci., 86, 881–890 (2002). 119. J.M. Duval, A.J.B. Kemperman, B. Folkers, M.H.V. Mulder, G. Desgrandchamps, C.A. Smolders, Preparation of zeolite filled glassy polymer membranes, J. Appl. Polym. Sci., 54, 409–418 (1994). 120. S. Husain, W.J. Koros, Mixed matrix hollow fiber membranes made with modified HSSZ-13 zeolite in polyetherimide polymer matrix for gas separation, J. Membr. Sci., 288, 195–207 (2007). 121. I.F.J. Vankelecom, S. Van den Broeck, E. Merckx, H. Geerts, P. Grobert, J.B. Uytterhoeven, Silylation to improve incorporation of zeolites in polyimide films, J. Phys. Chem., 100, 3753–3758 (1996).

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122. De Q. Vu, W.J. Koros, S.J. Miller, Mixed matrix membranes using carbon molecular sieves I. preparation and experimental results, J. Membr. Sci., 211, 311–334 (2003). 123. De Q. Vu, W.J. Koros, S.J. Miller, Mixed matrix membranes using carbon molecular sieves II. Modeling permeation behavior, J. Membr. Sci., 211, 335–348 (2003). 124. T.T. Moore, W.J. Koros, Non-ideal effects in organic-inorganic materials for gas separation membranes, J. Molecul. Struct., 739, 87–989 (2005).

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20 Nondestructive Testing of Composite Materials Zhongyi Zhang and Mel Richardson Advanced Polymer and Composites (APC) Research Group, Department of Mechanical and Design Engineering, University of Portsmouth, Portsmouth, UK

20.1 Introduction Nondestructive testing (NDT) is defined as testing techniques capable of identifying defects in materials and structures without causing any changes and damage. The location, geometry and severity of invisible damage are detected by identifying mechanical, thermal, optical, electrical, acoustic and radioactive abnormalities or irregularities associated with the components and structures under investigation [1–7]. NDT has been successfully employed in most stages of manufacture, service or maintenance of a component or structure in an attempt to:

r r r r r r r

Assist in product development Screen or sort incoming materials Monitor and improve manufacturing processes Verify proper processing Verify proper assembly Inspect in-service damage Validate the maintenance and repair tasks.

Composites typically consist of a continuous phase (matrix) and fibrous dispersed phases (reinforcement) with pronounced improvements in their properties and performance. Composites have been extensively

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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employed as load-bearing structural materials because they have many unique properties and performances. These properties and performances are dictated by parameters such as:

r r r r r r r

Reinforcing materials Matrix materials Interactions between reinforcing and matrix materials The ratio of fiberfiber to resin in the composites Orientation of fiberfiber reinforcement Size and dispersion of particle reinforcement Architectures of reinforcements.

Depending upon the matrices used, composites can be classified into metallic matrix composites (MMCs), ceramic matrix composites (CMCs) and polymeric matrix composites (PMCs). Polymer composites possess distinct properties such as high specific modulus and strength, light weight, high productivity, excellent processability, environmental degradation resistance and cost effectiveness. This chapter will focus on the principles and practical applications of NDT techniques which can be employed for damage evaluation in polymer composite components and structures. Polymer composites have sustained rapid development for many decades and found ever-increasing applications as engineering components and structures in the fields of aviation, aerospace, defence, marine, construction, transportation, energy, sports and leisure industries. Both Airbus A380 and Boeing’s 787 Dreamliner use large quantities of polymer composite components and structures for fuel efficiency and low operating cost. The proportion of polymer composites in civil and military aircraft is still steadily increasing. In these applications, it is critically important to evaluate the integrity and reliability of components and structures in a nondestructive way because routinely employed mechanical property testing is destructive in nature and not practically and economically viable in most circumstances. Failure in identifying potentially dangerous defects could possibly lead to malfunctions, failures or even catastrophic disasters. We have already witnessed many disasters which are directly related to an inability in detecting defects in polymer composite components and structures. Examples are train derailments, ship capsizing, building collapsing and aeroplane crashes. This emphasizes the need for reliable and effective implementation of NDT techniques. As far as polymer composites are concerned, defects can be introduced in many different ways during both manufacture and in service. There are many parameters involved in manufacturing processes which could result in the formation of defects. First of all, different reinforcements and matrices are combined together to achieve the required performance for components and structures. Secondly, the chemical reactions associated with thermosetting composites can lead to the generation of heat and the liberation of low molecular weight substances. Thirdly, defects can be introduced at different stages during fabrication processes. This is further complicated by thickness differences and the presence of slots, holes and inserts through which components can be assembled into a complete structure. Finally, damage is more likely to be introduced into polymer composites when they are subjected to different kinds of mechanical loading. They can be mechanically degraded to varying extents by the static, impact and cyclic loading conditions. Also, environmental factors can have an adverse effect, namely high temperatures, rapid temperature variations, humidity, ultraviolet radiation and high energy particle bombardment. Polymer composites exhibit complex damage mechanisms under different loading conditions because of anisotropic strength and stiffness. Instead of a predominant single crack often observed in most isotropic brittle materials, extensive damage throughout the specimen usually develops under loading. At a macroscopic level, basic failure mechanisms in polymer matrix composites include plastic deformation, matrix cracking, matrixfiber debonding, delamination and fiber breakage. Damage varies widely depending upon material properties, reinforcement stacking sequence and the nature of loading. Any combination of these failure mechanisms

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can lead to damage occurrence, resulting in varying degrees of reduced strength and stiffness [8–11]. In order to efficiently and reliably detect and identify the defects in polymer composite components and structures, it is necessary to develop a good understanding of the commonly-occurring failure mechanisms.

20.2 Failure Mechanisms in Polymer Composites Polymer composites fail by different mechanisms compared to metallic and ceramic composites due to their fundamental differences in composition, structure and manufacturing processes. Commonly-occurring failure mechanisms in polymer matrix composites include plastic deformation, matrix cracking, matrix/ fiber debonding, delamination and fiber breakage. 20.2.1

Matrix Deformation

Plastic deformation of a matrix is rarely considered as a failure. It is, however, an important feature mainly responsible for the greater toughness and impact resistance exhibited by many advanced polymer composites. It has been shown that matrices with high strains to failure offer excellent post-impact compression properties. The amount of shear deformation during failure can be determined by examining the fracture surface of the composites. The fracture surface would indicate if the matrix has been considerably drawn, which is a process involving a significant dissipation of energy. Shear flow in polymer composites is highly desirable because it can blunt the sharp cracks and result in a redistribution of loading fields. It has been demonstrated that specimens with a greater ability to undergo plastic flow exhibit greater fracture toughness with respect to crack initiation. Plastic deformation can give surface quality cosmetic problems but not necessarily have a significant effect on the structural integrity and load-bearing capabilities. There might not be immediate reduction in strength and stiffness. More and more academic and industrial interests are devoted to thermoplastic matrix composites because they exhibit excellent toughness and impact resistance mainly due to their extensive plastic deformation. 20.2.2

Fiber–matrix Debonding

In polymer composites, good adhesion between fiber and matrix is achieved via judicious formulation and careful control of fabrication parameters. Generally speaking, the fiber fracture strain is greater than that of the matrix. At low loading level, a small crack will be initiated at a point of stress concentration in the matrix. This crack is either halted by the fiber or will pass around the fiber without destroying the interfacial bond. When the applied load is increased, fiber and matrix deform differentially and a high level of shear stress is developed at the interface. When this shear stress exceeds the static interfacial shear strength, interfacial debonding occurs and then extends some distance along the fiber. It has been recognized that fiber–matrix debonding represents a localized failure mode that is usually difficult to detect using conventional techniques. The amount of debonding within polymer matrix composites depends upon the level of the surface treatment applied to the fibers during the manufacturing process. Generally, fibers with a low level of surface treatment will have a tendency to debond more easily because of the poor interfacial adhesion between the reinforcement and matrix. It is a common practice in industry that reinforcements are either chemically or physically treated to facilitate and enhance the bonding between the reinforcements and matrices. 20.2.3

Matrix Cracking

Matrix cracking is another significant type of damage and a common failure mode. Matrix cracking is commonly localizsed and difficult to detect. With increasing load, the density of cracks increases and appears

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to stabilize at a unique value for a given laminate. Although matrix cracking will not lead to significant reductions in mechanical performance, it can act as a precursor to delamination which is regarded as a detrimental mode of damage. In the case of impact loading on a flexible target, matrix cracking can initiate at the lower surface of the specimen and propagate upwards through the laminate. 20.2.4

Delamination

Delamination has been recognized as one of the most significant failure mechanisms exhibited by polymer composites. Delamination refers to the ply separation in a laminate. The skin/core separation in a sandwich composite is usually referred to as debonding but is also addressed as delamination due to the similarity. Delamination is the most difficult damage mode to deal with quantitatively because it only occurs between the plies but leave no traces on the surface. The invisibility is misleading and problematic which creates great difficulties in reliable and practical damage identification. The growth of delamination under monotonic or cyclic loading may result in a drastic reduction of strength and stiffness of polymer composite components and structures. The strength and stiffness reduction due to delamination varies depending on the delaminated area and type of loading for a given laminate. The driving force for delamination is the mismatch in elastic constants between adjacent plies that result in high interlaminar stresses. Although delamination is normally initiated at a free edge, it can also occur at intersections of matrix cracks in adjacent plies. Low velocity impact-induced damage is a typical example. The incident energy is not high enough to penetrate or fracture the materials or to create any visible defects on the surface but it will cause ply separation. Small areas of delamination are likely to lead to considerable reductions in the mechanical properties of such materials. 20.2.5

Fiber Breakage

Since fibers are the principal loading-bearing constituent in polymer matrix composite materials, fiber breakage can have an adverse effect upon both strength and stiffness. Fiber breakage in composites may occur for a number of reasons. Transverse impact loading often leads to the formation of localized fiber breakage at the beginning of impact. It has been suggested that fiber breakage can be the most detrimental factor in lowering tensile strength of polymer composite components and structures. Continuity and different textile structures of fiber woven fabrics play a crucial role in materialising physical and mechanical property enhancements. Loose fiber ends caused by fiber breakage are invariably the sites for the initiation and growth of fiber–matrix debonding due to the high shear stresses at the interface near the end of a fiber. It has been recognized that the energy associated with fiber breakage is considerably greater than that for matrix-dominated failure mechanisms. Consequently energy can be dissipated and absorbed in small areas. Fiber breakage is the principal energy-absorbing mechanism associated with complete penetration impact tests on polymer matrix composite materials. 20.2.6

Combination of Different Failure Modes

Commonly-occurring damage in polymer composites usually involves concurrent development of several different categories of the aforementioned damage. Generally, different types of damage have a tendency to be present and interact in complex ways. This leads to difficulties in understanding the characteristics of the damage development process. This is further complicated due to the fact that polymer composites are heterogeneous and anisotropic to varying degrees. Linear elastic fracture mechanics developed for homogeneous and isotropic materials is not readily applicable to polymer composites. It should be noted that not all failure

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mechanisms are necessarily present in a damaged component simultaneously. In most circumstances, one of them may contribute predominantly to the failure. The damage in an impacted specimen is a typical example in which different kinds of failure modes are readily recognizable. An indentation on front surface is an indication of plastic deformation. A crack on the back surface is associated with the matrix cracking, matrix/ fiber debonding and fiber breakage. Delamination is internally developed. These categories of damage in combination will contribute to significant reductions in strength and stiffness in polymer composite components and structures.

20.3 Visual Inspection Visual inspection is an NDT technique by which polymer composite components and structures are visually evaluated. It is known to be one of the oldest NDT techniques. In ancient times, artefacts made from different materials were visually inspected by unaided eyes and abnormalities determined. Visual inspection continues to be a viable technique due to the fact that it is simple, speedy and cost effective compared to other sophisticated NDT methods. Whenever NDT is required for polymer composite components and structures, visual inspection should invariably be the first choice. The suspicious artefact is typically visually inspected before sophisticated NDT techniques are considered. For example, surface finish degradation and indentation on a surface can indicate the presence of possible internal damage and the location for further investigation. Visual inspection has been successfully used for metallic and ceramic composites. Unfortunately, it is not equally applicable to polymer composites due to the fact that they are fabricated in different ways. Such composites have a tendency to fail by different mechanisms from metallic and ceramic composites when subjected to similar loading conditions. When these materials are subjected to low and intermediate level stress, damage is usually initiated and developed internally in the form of delamination, leaving little or no damage indication on the surface. The primary disadvantage of visual inspection is that the components can only be superficially examined. It cannot generate complete and accurate information of the damage. A typical example is, when polymer composites are subject to low velocity impact loading, there is barely visible damage or only a tiny indentation or crack. Actually extensive damage may have been developed internally in the form of delaminations and the structural integrity has already degraded and load-bearing capability substantially reduced. There is another factor that makes visual inspection less attractive and practical. In the case of evaluating a component with simple geometry, visual inspection of both front and back surfaces can be conveniently carried out because they can be practically accessed. However, it is often impossible to examine back surfaces of polymer composite components and structures in most circumstances due to restricted physical accessibility. Typical examples are the difficulties in accessing to the inner surfaces of different pipes, cylinders, containers and geometrically complex components and structures. Consequently, advanced NDT techniques are expected to be capable of identifying the damage from front surfaces where there are usually no physical accessibility restrictions. Therefore, specific procedures and personnel training are necessary to guarantee the success of an evaluation. Visual inspection must be carefully implemented and controlled in an attempt to obtain accurate and reliable results. Visual inspection alone cannot be single-handedly employed to carry out a conclusive evaluation. Instead, it should be used in conjunction with other sophisticated NDT techniques. Techniques to enhance human visual inspection continue to be developed, such as optical imagery, fiber optics, magnification and computer-aided diagnosis of defects. Liquid penetration testing is a variation and extension of visual inspection. Cracks, porosity and discontinuities open to the surface can be readily detected. The color penetrant is drawn into very small openings by a capillary effect when it is applied to a surface. The color penetrant is normally left on the surface for a specific period of time for permeation and penetration and

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then excessive liquid is removed from surface. When the article is viewed under illumination, the damage details are revealed with better visibility and improved accuracy due to color differences.

20.4 Acoustic Emission Acoustic emission is an NDT technique by which polymer composite components and structures are acoustically evaluated as illustrated in Figure 20.1. It is also called stress wave emission. It involves detection of acoustic signals that are spontaneously generated by materials when they undergo progressive deformation and failure in different loading conditions. A set of sensors is judiciously attached to different locations according to the evaluation scheme, and both audible and inaudible acoustic signals are collected for analysis and interpretation. In general, acoustic emission can be used to monitor the variation in material conditions in real time and to determine the location of emission centers as well [12–15]. For example, when a piece of wood or bamboo stick is flexurally stressed, audible noises can be heard, indicating the crack initiation, development and catastrophic failure. Since it is usually a passive technique, no equipment is required to produce a pulse for defect detection. It is highly sensitive to growing cracks and delaminations. Additional advantages are its ability to monitor the entire component at the same time. It has been extensively and successfully employed in testing composite pressure vessels with significant success. With remote monitoring it can be used in hostile environments. Furthermore, the component under study can usually remain in operating condition so evaluation can be implemented at reasonable cost. It has been proved that acoustic emission is an ideal alternative to other long-term, in-service monitoring techniques. The primary applications of acoustic emission in polymer composites are concerned with the assessment of damage status, failure mechanisms and failure locations. Real-time monitoring can be carried out on specimens subjected to different loading conditions such as tensile, compressive, flexural, fatigue and thermal differential stressing. Acoustic emission signal parameters such as acoustic emission event counting, amplitude distributions, frequency distribution and duration time are analyzed and correlated to the damage because different failure mechanisms give rise to unique acoustic events. In general, low, intermediate and high

Preamplifier

Filters

Sensors

Amplitude (V)

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Loading

Printer or Recorder

Figure 20.1

Schematic of acoustic emission.

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amplitude distributions can be correlated to matrix cracking, delaminations and fiber breakage, respectively. Similar correlation can be established by the real-time frequency analysis of acoustic emission signals. Unfortunately, acoustic emission suffers from some inherent limitations. Monitoring systems are passive in nature. They simply listen to sounds generated by defect initiation and their development in the stressed state. It cannot detect existing defects because no audible signals are generated after the defects have developed. It is always difficult to accurately determine defect details because not all materials emit sound waves from defects under load and the propagation of sound may be dampened depending upon the nature of the materials. The correlation of acoustic signals to the defects is sometimes problematic and subjective to some extent. In critical circumstances, it is necessary to supplement testing results with other NDT techniques. It is also important to take the component geometry and possible sensor placements into consideration in implementing damage evaluation. Large numbers of sensors mean good testing accuracy but at high cost. Furthermore, the great sensitivity of the monitoring system might impose additional problems. Electrical interference and ambient noise must be filtered out of emission signals and the multiple channels of travel paths from the source to the sensor in complex structures can make signal identification and interpretation difficult.

20.5 Ultrasonic Scanning Ultrasonic scanning is an NDT technique by which polymer composite components and structures are also acoustically evaluated as illustrated in Figure 20.2. However, it is significantly different from acoustic emission. It is accomplished by introducing electronically controlled pulses into a material via its outer surface. The ultrasonic energy then travels within the material and finally reaches an outer boundary. Material conditions are diagnosed from the characteristics of received or returned ultrasonic signals [16]. NDT can be accomplished either by pulse-echo inspection or through-transmission inspection depending upon the specific requirements. Pulse-echo is a technique where a pulsed ultrasonic beam is transmitted

α: Start Pulse β: Defect Echo γ: Bottom Echo

Scanning Device

α

β

Y X

Probe Water

Water Tank

Figure 20.2

Schematic of ultrasonic C-scan.

γ

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through a coupling agent into the material under study and travels to another surface within the material. The pulse is then reflected back to a transducer that may or may not be the transmitting transducer. Flaws in the material will reflect or scatter part of the incident energy. Therefore, a flaw echo will appear on the oscilloscope at the location corresponding to flaw depth in the material. Through-transmission is used in several cases, particularly for highly acoustically attenuative material where a pulse echo suffers significant loss in signal amplitude. The instrumentation used in pulse-echo testing can be readily employed in through-transmission testing. Three scan modes are implemented in ultrasonic scanning, referred to as A, B and C-scans, respectively. A-scan produces single point information concerning the object quality. The signal amplitude and its location relative to those of the signals corresponding to the upper and lower surfaces of the object give an indication of the severity and the through-thickness location of the damage. B-scan is essentially a linear collection of A-scan signals that are equivalent to scanning a slice through the sample. The third is ultrasonic C-scan that has been proved to be the most popular and useful of all scanning modes. In C-scan the component is typically immersed in a water bath. The transducer sweeps backwards and forwards across the component, receiving and analyzing the signal reflected from different surfaces of the damaged specimen. The data is analyzed and the signal amplitude is converted into a voltage. Finally, the data is visualized in terms of differences in amplitude attenuations and time-of-flight signals for interpretation and evaluation. Damage evaluation of a low velocity impacted polymer composite panel using ultrasonic C-scan is shown in Figure 20.3. Ultrasonic C-scan technique is ideally applicable to the detection of certain kinds of defects. As far as NDT of polymer composites is concerned, ultrasonic C-scan is preferentially employed because it can produce fairly satisfactory results in comparison with other conventional NDT techniques. It has been successfully used for investigating failure mechanisms and for identifying the damage in polymer composites. Although ultrasonic scanning has been demonstrated to be a successful NDT technique for polymer matrix composite materials, it suffers from some inherent disadvantages. To a great extent, the successful damage evaluation depends upon the coupling of the ultrasonic transducer with the component or structure under

Figure 20.3 Damage visualization by ultrasonic C-scan for a GRP panel subjected to low velocity impact by a hemi-spherical impactor.

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investigation. Ultrasonic waves can be drastically attenuated in air, especially at high testing frequencies. Furthermore, differences in acoustic impedance of air and solid materials cause most of the ultrasonic energy to be reflected at the surface of the solid rather than to be propagated in it. Consequently, a liquid coupling agent or direct transducer contact is necessary to introduce ultrasonic energy into solid materials. In most circumstances, direct transducer contact is practically impossible due to the complicated geometries of the components. The liquid coupling agent can be water in a tank where the object is immersed. A water jet ejected from a nozzle is used to provide an acoustic path to test large components and structures as well as honeycomb structures because they would float on top of water as a result of their lower density. It should be noted that excessive acoustic attenuation of ultrasonic waves in some materials can lead to damage evaluation failure. It is important to understand the responses of materials to ultrasonic waves prior to evaluation. This problem is further compounded due to the reality that most polymer composites components and structures have complex geometry. In addition, scanning is not a whole field and real-time testing technique. It can be tedious and time consuming to scan large surface areas. The ultrasonic C-scan technique may not be able to detect small-scale fiber fracture, matrix cracking or micro-mechanical damage mechanisms. In addition, the ultrasonic C-scan technique produces only a two-dimensional view of the defect or damage zone, giving no through-thickness data. More information can be obtained from a time-of-flight analysis to generate a three-dimensional presentation of the damage zone. In critical circumstances, it is recommended that the ultrasonic scanning technique be employed in conjunction with other NDT techniques or destructive testing techniques in an attempt to achieve a damage evaluation with high acceptability and reliability.

20.6 Radiography Radiography is an NDT technique by which polymer composite components and structures are radioactively evaluated as illustrated in Figure 20.4. Radiography is used for both medical and industrial applications based on the same principles. Differences in absorptions or scattering of X-rays are induced when they pass through

Radiation Source

Digital Camera

Defective Sample

Recording Plate

Figure 20.4

Schematic of radiography for NDT.

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the damaged and undamaged areas. Flaws will either allow more or less X-ray to pass, absorb and scatter. These significant differences can be used to indicate the presence of damage [17]. Radiography has been widely employed to evaluate the structural integrity of polymer matrix composites nondestructively. It is capable of identifying damage such as matrix cracking and delamination as well as extensive fiber fracture. Penetration enhanced in situ radiography enables real-time analysis of the damage development, deformation mechanisms and material failure. It has been demonstrated that the radiographic inspection can reveal useful information about voids, entrapped foreign materials, matrix cracks, resin-rich and resin-starved areas. Radiography is particularly useful for the detection of honeycomb core defects in bonded sandwich assemblies. Low density and thin composite skins usually provide minimal interference to X-rays when imaging honeycomb core materials. Core defects such as blown, crushed, condensed, fatigued, and corroded areas can be readily detected. It is also capable of detecting water intrusion into the honeycomb. X-ray computed tomography (CT) was originally developed in medicine for visualization of soft tissues and bones. It was subsequently adapted and extended to a wide variety of material science applications. CT can provide nondestructively three-dimensional maps of X-ray attenuation within different materials. The basic CT configuration consists of an X-ray source, a sample through which the X-ray can pass and a detector assembly, which measures the attenuation as the X-ray passes. Many different pathways are measured by either rotating the sample or the source-detector assembly. Subsequent acquisition of many such slices results in a three-dimensional image of X-ray attenuations which provide detailed information on damage location, size, geometry and depth. Like all techniques, radiography has certain disadvantages. First of all, the accessibility to the opposite surface of an object under investigation is always required where radiation variations are detected and imaged. Secondly, radiation sources are expensive to generate, maintain and operate. X-ray radiations are health hazards. It is, therefore, necessary to close off the area when making exposures in order to avoid endangering the public. Personnel who operate radiography must be qualified to work with radioactive sources. Finally, identification of various failure modes in polymer composites can be difficult. It can indicate the presence of defects but provide little information about the failure mechanisms involved. It is recommended to employ radiography in combination with other NDT techniques in order to make a reliable damage evaluation.

20.7 Thermography Thermography is an NDT technique by which polymer composite components and structures are thermally evaluated, as illustrated in Figure 20.5. It includes all methods which use heat sensing devices or substances to measure the resultant temperatures and thermal gradients of an object under inspection when it is heated up or cooled down. The infrared camera is the most important sensing device and has been successfully used for decades. Advances in optical and electrical technology offer infrared cameras with high sensitivity and make it possible to differentiate minute temperature differences. The damage is visualized by the infrared camera in the form of hot spots due to the differential absorption and dissipation of heat in damaged and undamaged areas [18]. Thermography is a whole field, noncontacting, safe and rapid technique which is typically employed to evaluate the structural integrity of different polymer composites. Delaminations, blind surface impact damage and surface cracks can be detected by infrared thermal field techniques. Thermal sources are employed to excite the defects in materials under inspection. They can be either actively or passively implemented. Two categories of methods are used to generate thermal fields for damage detection. The first is referred to as externally applied thermal excitationin, in which an external heat source is applied to the surfaces in order to generate thermal gradients in the component or structure. Devices such as hot air guns, flash lights, ovens and refrigerators have been used for this purpose. In the absence of flaws,

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Thermal Camera

Externally applied thermal excitation

Figure 20.5

Cyclic loading generated thermal excitation

Schematic of thermography for NDT.

unperturbed isotherm patterns indicate good uniformity dictated by thermal exchanges with environment. The patterns can be different depending upon what the heat device is and how it is employed. In the case of the presence of a flaw or a damaged region, the thermal absorption behavior upon heating and then the thermal exchanges with the ambient environment will lead to irregular thermal conductivity which causes a perturbation in the heat flux and temperature field. Degraded uniformities and abnormalities are used to evaluate the damage. The second method is called cyclic loading generated thermal excitation. Monotonic loading such as tension or compression is not an appropriate way to thermally excite the defect because it generates limited heat and hot spots cannot be introduced. Repetitive loading/unloading below the yield and fracture points of materials or other cyclic loading will generate heat internally for damage evaluation. Although most materials can be considered as linear elastic over a significant portion of their loading range, there is always a hysteresis loop when subjected to loading and unloading due to the elastic–plastic nature of materials, and this is particularly true for polymer matrix composites. The energy generated in a cycle of loading and unloading is dissipated in the form of heat. The greater the magnitude of stress and deformation, the more generated heat there is. Since cracks, flaws and damaged areas will act as stress concentrators under applied stress, the heat generated in cycles of loading and unloading will be disproportionably high in the damaged area. Polymer composites exhibit low thermal conductivity so thermal abnormalities will exist for a while, leading to the formation of hot spots in the region immediately adjacent to the flaws.

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Like other NDT techniques, thermography is far from being perfect and has some disadvantages. The effectiveness of thermography depends upon the thermal conductivity of the materials under examination. If the object surface contains different materials, it will present a range of surface thermal emissivity. This can lead to difficulties in interpretation. Surfaces with low emissivity do not emit enough radiation for adequate measurements so it is necessary to apply coatings to them. Extraneous background thermal noises have an adverse effect on damage evaluation. It is necessary to improve software and functionalities of thermography in order to distinguish small temperature differentials and to filter out thermal background for improved damage detection capability. It should be noted that thermal abnormalities identified by thermography do not necessarily mean there are defects which would sacrifice the structural integrity of the polymer composites. Therefore this technique should be employed in conjunction with other NDT techniques which are able to mechanically determine irregularities and abnormalities.

20.8 Laser Interferometry Laser interferometry is an NDT technique by which polymer composite components and structures are mechanically evaluated. Both holographic interferometry (HI) and electronic speckle pattern interferometry (ESPI) are in the category of laser interferometry. HI has been successfully employed to surface displacement measurements, contouring and vibration mode studies of different components and structures. HI has already been extended to the NDT community with some success [19]. In theory, defects usually induce anomalies in displacement fields and vibration modes. They can be readily diagnosed by interpreting the perturbations and interruptions of interferograms. Many attempts have been made to apply HI to the evaluation of damage in a nondestructive way. The suitability and applicability of HI as an NDT technique has been experimentally demonstrated. Its primary applications are the characterisation of laminate structures, composite materials, cylinder bores, turbine blades, and pneumatic tyres as well as solid propellers. The main handicap associated with HI is the stringent stability requirement of the operational environment. The extraordinary high interferometric sensitivity to displacement imposes a limitation that the components under investigation and the surrounding environment must be extremely stable otherwise the fringe patterns will de-correlate. This is further complicated by the fact that the recording time is a function of the recording medium sensitivity, component size and laser power. The required exposure time is usually in the order of seconds or tens of seconds for holograms. In addition, the chemical photographic processing of holograms is laborious, time consuming and cumbersome. Therefore, HI has been substantially confined to laboratories and found limited industrial applications. ESPI has been developed to overcome the difficulties experienced by HI. The principle and configuration of ESPI are illustrated Figure 20.6. In ESPI, traditional time-consuming photographic film recording is replaced by a video camera. Consequently, experimental results can be accessed in real time. The combination of laser, video techniques and electronic and digital signal processing systems means that ESPI offers great potential for rapid data acquisition and automatic analysis in engineering measurement and testing with high sensitivity to displacement and deformation. Recently ESPI has been recognized as one of the most promising laser interferometric metrologies due to the fact that it is capable of making full-field, non-contacting and real-time measurements. There has been considerable success in various applications, and damage evaluation of materials is one of its important applications. The geometrical features and magnitudes of internal damage can be determined in terms of abnormalities and interruptions of fringe patterns which can be observed on a TV monitor in real time as shown in Figure 20.7. The internal damage can be visualized in 3-D format as shown in Figure 20.8. The formation of these patterns is due to the fact that damage reduces the strength and stiffness of materials, leading to different loading responses between damaged and undamaged areas. Undamaged areas have a tendency to be deformed

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Mirror

Laser Mirror

Beam Splitter

Lens Lens

Digital Correlation

Camera

Figure 20.6

Sample

Schematic of electronic speckle pattern interferometry (ESPI).

or displaced less than the damaged area when subjected to the same loading condition. A boundary between undamaged and defective areas is likely to develop due to differences in stiffness. The lower stiffness of the damaged areas is responsible for the higher density and abnormality of fringe patterns due to excessive deformation in this region. Unlike thermal excitation, the fringe patterns triggered by mechanical excitation can stabilize for a long period due to the permanent and irreversible deformation associated with the damaged areas. In some circumstances, the displacement differences may be very small in magnitude, but ESPI is capable of differentiating between them due to its high sensitivity [20–22]. As far as damage evaluation using both HI and ESPI is concerned, excitation techniques play an important part in successful defect detection. Four types of excitation techniques are commonly employed, namely

Figure 20.7

Irregularities in fringe patterns showing a defect in a polymer composite panel.

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1.85 1.27 0.70

Z

0.13 –0.44 –1.01 –47.28

53.63

–27.17

33.90

–7.05

14.18 X

13.06

–5.55 –25.27

Y

33.18

–45.00 53.29

Figure 20.8

3-D visualization of internal damage by ESPI.

mechanical stressing, thermal radiation, pressure difference and vibration. Judicious selection of excitation technique is very important in an attempt to generate practical results at low cost. Although the requirement for environmental stability is greatly relaxed for ESPI compared to HI, it still imposes restrictions on its wider acceptance in many applications. Fringe patterns can be registered by rigid body movement which might lead to some problems in results analysis and interpretation. Extra care is needed in interpreting the damage in terms of abnormalities of correlation fringe patterns. The determination of geometrical features and magnitudes of internal damage might be difficult for inexperienced users. Optical phase extraction can be used to filter out the background speckle noises but it requires additional hardware and software for data acquisition and processing.

20.9 Electronic Shearography Electronic shearography (ES) is also called electronic speckle pattern shearing interferometry (ESPSI), or simply shearography. The principle and configuration of ES are illustrated in Figure 20.9. Although ES is another laser interferometric technique, it has been extensively used for damage evaluation of polymer composites with significant success. ES gets its name from the shearing characteristics of the camera used. ES was originally developed to measure object surface strain distributions using a special configuration capable of optically differentiating the surface displacement fields. There are two types of optical arrangement for ES. The first is based upon a skewed Michelson interferometer configuration in which the object is illuminated by a single wave front inclined at a small angle and viewed through a Michelson interferometer. It has great flexibility to change the shearing direction and magnitude but with low laser illumination efficiency. The second type of ES employs a thin transparent glass wedge to cover half of the lens aperture. This optical

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Laser IIIumination

Shearing Component

Lens Aperture

Digital correlation

Figure 20.9

Schematic of electronic shearography(ES).

configuration has good illumination efficiency but low flexibility. Both optical configurations have been successfully employed [23, 24]. ES has been recognized as a viable NDT technique due to the fact that it is capable of making a full field, non-contacting and real-time measurements of displacement derivative fields. Unlike ESPI, the fringe patterns represent surface displacement derivatives with respect to the shearing direction instead of the displacements. In theory, surface, subsurface and interior defects usually cause strain concentrations around the periphery of defects. Since strain concentrations induce perturbation and irregularity on the surface displacement derivative fields, the anomalies in a shearogram can be identified and correlated with defects inside the materials. Typical shearogram is shown in Figure 20.10. ES measures the first derivatives of displacements in the shearing direction as illustrated in Figure 20.11 so rigid body motion does not produce any strain. Thus strain concentrations associated with defects can be easily revealed. These further relax the requirement of environmental stability and create the opportunity for ES being used in real industrial environments. Portable ES testing systems for NDT are now commercially available in which all optical and electronic devices are integrated into a vacuum hood. The vacuum is applied to seal off the testing area and to excite the damage. The damaged area will be drawn toward the hood more than the undamaged area so it is readily visible. Internal damage evaluation of a sandwich decking panel for a marine yacht is shown in Figure 20.12. Like all other NDT techniques, ES is by no means perfect. It is difficult to visualize and quantify defects because damage evaluation by ES is on the basis of anomalies in the displacement gradient fields. The damage visualization is distorted by optical shearing to some extent, depending upon the optical configurations.

20.10 Optical Deformation and Strain Measurement System Optical deformation and strain measurement is based upon the fact that the distribution of gray scale values of a rectangular area in the undeformed state corresponds to the distribution of gray scale values of the same area in the deformed state. The principle and configuration is shown in Figure 20.13. It combines the advantages of photogrammetry and the object grating method.

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0.58 16.7 0.44 0.31 0.18

44.0 72.2

0.04

99.6

–0.09

128.2

–0.22

155.6

–0.36

184.2

–0.49 –0.63 22.4

Figure 20.10

46.2

69.0

92.8

115.5 139.0 162.1 185.6 208.3 232.1

Typical shearogram indicating the presence of internal damage in a polymer composite panel.

Photogrammetry is one of the optical methods that lead to 3D-coordinates of surface points. The displacement vectors, local strain values and contour difference can be computed from the data when the object is deformed. When the object points on the surface of the specimen are arranged like a grating, this is the technique well known as the grating method in experimental mechanics. Instead of an expendable line mesh, a stochastic pattern is applied to the surface using a graphite spray which allows a high resolution. The optical deformation and strain measurement system only require ordinary light illumination and possess many unique features including simple sample preparation, large measuring area, noncontact, full field, material independent determination and three-dimensional visualization [25]. The optical arrangement in optical deformation and strain measurement consists of two CCD cameras and a loading device. If the position of the two cameras and two homologous image points is known, the corresponding object point can be calculated. This procedure is known as space intersection. Therefore, a geometric model must be established, which implements the transformation from image points to object points. The parameters of the camera are necessary to establish the model that is calculated through a calibration procedure. The calculation of homologous image points of two deformation states recognized from two cameras can be achieved by a combination of 2D displacement calculation processes. It is important that one image is defined as the reference state. After determining all 2D displacement fields, the homologous image points can be calculated. With this information, the object coordinates for each state can be calculated by space intersection. After the determination of the displacement field, the strain distribution can be numerically calculated. The

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0.58

0.34

0.10

Z –0.14

–0.38

–0.63 –39.24 16.62 72.47 128.33 Y 184.19 240.04 –0.20

46.27

139.20

92.73

185.67

232.13

X

Figure 20.11

3-D visualisation of internal damage by ES in terms of the first derivatives of displacements.

Figure 20.12

Inspection of delaminations in a composite decking panel by a portable ES system.

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yt

y

Displacement Determination

xt

x

Deformed

Un-deformed

Strain Determination

z

x Grating

Figure 20.13

y

Schematic of optical deformation and strain measurement system.

way to calculate the surface strain is through the transformation of the 3D displacement distribution into a 2D displacement distribution. Finally, the strain can be calculated in the 2D spaces according to the algorithms. Damage identification for polymer composites using the optical deformation and strain measurement technique is based upon determining the irregularities and abnormalities of deformation and strain fields. Mechanical loading is normally progressively applied to the component and images are grabbed from each loading levels. Both external and internal defects in polymer composites would weaken the materials, leading to the reduction in strength and stiffness. The damaged area would be deformed and strained more than the undamaged area so the defects can be visualized in terms of deformation and strain differences. Visualization of impact-induced, barely visible damage in polymer composites using an optical deformation and strain measurement system is shown in Figure 20.14.

Figure 20.14 Visualisation of an impact induced barely visible damage in polymer composites using an optical deformation and strain measurement system.

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The principle of NDT for polymer composites using an optical deformation and strain measurement system is similar to that in ESPI and ES. All of them use optical devices and configurations to mechanically evaluate materials directly associated with structural integrity but it is different from ESPI and ES in other ways. No laser illumination and correlation are involved in the optical deformation and strain measurement system so it eliminates the requirement for environmental stability. However, its sensitivity is lower than that of ESPI and ES so higher loading levels are required to induce enough deformation and strain for damage identification. In other words, defects can be readily identified in less stiff and less strong materials. Therefore, laser interferometry including ESPI and ES can be complementarily employed with optical deformation and strain measurement system to nondestructively evaluate a wide variety of polymer composite components and structures.

20.11 Summary Structural integrity of polymer composites can be nondestructively assessed by evaluating their optical, thermal, acoustic, radioactive and mechanical properties and attributes. Some other NDT techniques which are commonly practiced to evaluate metallic and metallic alloys are not discussed due to the fact they are inappropriate to polymer composites. It should be noted that all NDT techniques have their own advantages but none are perfect. They all suffer from some inherent limitations and disadvantages. Successful damage evaluation and identification can be an expensive and arduous task. In general, each individual NDT technique is usually employed in conjunction with other NDT techniques in order to achieve high damage evaluation accountability and reliability especially when structural integrity is critically important such as for aviation and aerospace applications. It is extremely important to improve existing NDT techniques and to develop novel ones in order to keep abreast with ever-increasing applications of polymer composites as primary and secondary structures in which structural integrity and reliability are crucial and vital.

References 1. C. Boller, F. K. Chang, Y. Fujino, Encyclopaedia of Structural Health Monitoring, John Wiley & Sons Inc., 2009. 2. R. Jones and C. Wykes, Holographic and Speckle Interferometry, Cambridge University Press, Cambridge, 1989. 3. M. O. W. Richardson and M. J. Wisheart, Review of low-velocity impact properties of composite materials, Composites Part A: Applied Science and Manufacturing, 27, 1123–1131 (1996). 4. R. D. Adams and P. Cawley, A review of defect types and non-destructive testing techniques for composites and bonded joints, NDT International, 21, 208–222 (1988). 5. E. Bayraktar, S.D. Antolovich and C. Bathias, New developments in non-destructive controls of the composite materials and applications in manufacturing engineering, Journal of Materials Processing Technology, 206, 30–44 (2008). 6. G. Scott and C. M. Scala, A review of non-destructive testing of composite materials, NDT International, 4, 75–86 (1982). 7. Y. Y. Hung, Y. S. Chen, S. P. Ng, L. Liu, Y. H. Huang, B.L. Luk, R. W. L. Ip, C. M. L. Wu, P. S. Chung, Review and comparison of shearography and active thermography for non-destructive evaluation, Materials Science and Engineering: R: Reports, 64(5–6),73–112 (2009). 8. I. M. De Rosa, C. Santulli, F. Sarasini, M. Valente, Post-impact damage characterization of hybrid configurations of jute/glass polyester laminates using acoustic emission and IR thermography, Composites Science and Technology, 69(7–8), 1142–1150 (2009). 9. A. Vieira, R. de Oliveira, O. Fraz˜ao, J.M. Baptista, A.T. Marques, Effect of the recoating and the length on fiber Bragg grating sensors embedded in polymer composites, Materials & Design, 30(5), 1818–1821 (2009).

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10. S-G Kang, D-H Kang, C-Gon Kim, Real-time monitoring of transverse thermal strain of carbon fibre reinforced composites under long-term space environment using fibre optic sensors, NDT & E International, 42(5), 361–368 (2009). 11. G. J. Williams, I. P. Bond, R. S. Trask, Compression after impact assessment of self-healing CFRP, Composites Part A: Applied Science and Manufacturing, 40(9), 1399–1406 (2009). 12. L. Dong and J. Mistry, Acoustic emission monitoring of composite cylinders, Acoustic emission monitoring of composite cylinders, Composite Structures, 40,43–53 (1997). 13. I. M. De Rosa, C. Santulli, F. Sarasini, Acoustic emission for monitoring the mechanical behaviour of natural fibre composites: A literature review, Composites Part A: Applied Science and Manufacturing, 40(9), 1456–1469 (2009). 14. L. P. Dickinson and N.H. Fletcher, Acoustic detection of invisible damage in aircraft composite panels, Applied Acoustics, 70,110–119 (2009). 15. L. P. Dickinson, N.H. Fletcher, Acoustic detection of invisible damage in aircraft composite panels, Applied Acoustics, 70(1), 110–119 (2009). 16. M. Kobayashi, C.-K. Jen, J.F. Bussiere, K.-T. Wu, High-temperature integrated and flexible ultrasonic transducers for non-destructive testing, NDT & E International, 42(2), 157–161 (2009). 17. P. J. Schilling, B. R. Karedla, A. K. Tatiparthi, M. A. Verges and P. D. Herrington. X-ray computed microtomography of internal damage in fibre reinforced polymer matrix composites, Composites Science and Technology, 65,2071–2078 (2005). 18. M. Genest, M. Martinez, N. Mrad, G. Renaud and A. Fahr, Pulsed thermography for non-destructive evaluation and damage growth monitoring of bonded repairs, Composite Structures, 88,112–120 (2009). 19. J. A. Leendertz, Interferometric displacement measurement on scattering surfaces utilising speckle effect, Journal of Physics E: Scientific Instrument, 3,214–218 (1970). 20. Z. Y. Zhang and M. O. W. Richardson. Low velocity impact induced damage evaluation and its effect on the residual flexural properties of pultruded GRP composites, Composites Structure, 81,195–201 (2007). 21. M. O. W. Richardson and Z. Y. Zhang, ESPI non-destructive testing of GRP composite materials containing impact damage, Composites Part A: Applied Science and Manufacturing, 29,721–729 (1998). 22. M. O. W. Richardson and Z. Y. Zhang, Application of phase stepping ESPI to non-destructive testing of GRP composite materials, International Conference of NDT in Civil Engineering, Journal of The British Institute of Non-destructive Testing, 40(3), 183–187 (1998). 23. Y. Y. Huang, Shearography: a new optical method for strain measurement and non-destructive testing, Optical Engineering, 21(5), 391–395 (1982). 24. Y. H. Huang, S. P. Ng, L. Liu, C. L. Li, Y. S. Chen, Y. Y. Hung, NDT&E using shearography with impulsive thermal stressing and clustering phase extraction, Optics and Lasers in Engineering, 47(7–8), 774–781 (2009). 25. Z. Y. Zhang and M. O. W. Richardson, Visualisation of Barely Visible Impact Damage in Polymer Matrix Composites Using Optical Deformation and Strain Measurement System, Composites A: Applied Science and Manufacturing, 36(8), 1073–1078 (2005).

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21 Ageing and Degradation of Multiphase Polymer Systems Xavier Colin PIMM (UMR 8006), Arts et M´etiers ParisTech, Paris, France

Gilbert Teyssedre Laplace, Universit´e Paul Sabatier, Toulouse, France

Magali Fois Universit´e Paris-Est, CERTES EA 3481 – Centre d’Etude et de Recherche en Thermique, Environnement et Syst`emes, Cr´eteil, France

21.1 Introduction 21.1.1

Issues Associated with Material Ageing

Ageing can be defined as a slow and irreversible variation as a function of time (in use conditions) of a material structure, morphology or composition leading to a detrimental change in its use properties. The cause of this change can be the own material instability or its interaction with the environment of exposure. The definition so given is that viewed from an application point of view. There are issues associated with this definition that deserve to be mentioned. First, there are ageing mechanisms, essentially of a physical nature as detailed below, that are not irreversible in nature (e.g. crystallinity change, structural recovery, water uptake without loss of the integrity of the chemical structure, etc.), but may lead to a change in the use properties of these materials. The reversibility is in principle achievable by, e.g. thermal treatment or drying. However, this is not necessarily compatible with the use of materials as pieces, or the material will evolve again anyway when exposed to use environmental stresses. Second, irreversible material evolution in itself does not necessarily imply a detrimental change of use properties. It can even result in an improvement of properties. This represents indeed a marked difficulty when attempting to define so-called ‘ageing markers’ for materials, i.e. material properties to be monitored for health monitoring purpose: the marker must be Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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sensitive enough so as to provide the early signs of material evolution but, at the same time, there should be a correlation between the evolution of the marker and the changes in use properties. In the following, we shall consider ageing from the point of view of molecular, macromolecular or microscopic changes induced on the polymer, with room given on the impact on use properties. A large reference to the impact of various ageing mechanisms on mechanical properties will be made, in part because of the area of expertise of the authorship, and most importantly because of the more advanced knowledge and earlier consideration of the relationships between mechanistic ageing processes and macroscopic mechanical properties. This is even more evident when dealing with polymer-based composite materials which have emerged for the huge improvement of mechanical properties brought about in respect to the matrix alone, and which do still represent the essential part of the quantity of material being produced and a significant part of the research being carried out on the topic. Any polymer or composite piece has a mechanical role, and even in demanding applications such as electrical insulation, it can be more advisable to use composites instead of polymers alone for mechanical or thermal criteria (along with economics), even though it can be significantly detrimental to the electrical performances. In another respect, the policy of switching to all-composite epoxy-carbon aircrafts has been guided by mechanical engineering and the perspective offered in weight reduction with, to our eyes, little consideration of the consequences it may have on electrical networks onboard and on lightning issues. The challenges when dealing with ageing of materials are multiple, going from mechanistic processes, to materials selection and to design. In terms of area of expertise, it implies gathering people with skills ranging from macromolecular chemistry to engineering resorting to the various area of application of materials. The main tasks that are currently addressed, irrespective of materials considered and targeted application, are as follows: (a) (b) (c) (d)

understand how the material is evolving find the ageing markers relevant to the expected service properties define ‘end-of-life’ criteria and forecast material evolution (life estimation) find ways to prevent material ageing.

21.1.2

Classification of Ageing Types

There is a wide variety of polymer ageing types. Each one is a complex phenomenon, but decomposable into a set of elementary mechanisms that can be interactive: chemical reactions, macromolecular motions, molecular species transport into the polymer, crystallinity change, etc. Unfortunately, there is no entirely satisfactory method for their classification, because the most rigorous criteria, in theory, are often open to criticism in practice. The simplest and most agreed method consists in distinguishing two main ageing families according to the nature of involved mechanisms:

r r

chemical ageing is responsible for a variation of the macromolecule chemical structure physical ageing affects the macromolecule conformation, the material morphology or composition, without altering the macromolecule chemical structure.

Then, the different ageing types, defined in terms of the main ageing driving environmental stress, involve one of the two families of mechanisms defined above:

r r

Thermal ageing is initiated by material’s own instability. Its kinetics essentially depends on temperature and atmosphere composition. It can be a chemical or a physical ageing, or a combination of them. Humid ageing is due to a chemical or a physical interaction of polymer with water, the latter being either in a liquid or a vapor state.

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r r r r r

799

Photo-chemical ageing results from polymer interaction with light rays (in general, UV solar radiations). Atmospheric oxygen plays, in general, an important role (photo-oxidation). Radio-chemical ageing is due to the polymer interaction with ionizing (γ , β, α, neutrons, etc.) rays. Atmospheric oxygen can play an important role, especially at very low dose rate (radio-oxidation). Electrical ageing due to the interaction of electrical charges with the polymer or to a direct field effect on it. Mechanical ageing through, e.g. fatigue exerted on the polymer which can induce chain breaking along with morphological changes within the material. Biochemical ageing is due to the polymer interaction with living organisms (bacteria, molds, etc.).

Ageing is a difficult problem to study in practice, because it usually proceeds slowly in (soft) service conditions of materials and lifetime reaches typically several dozens of years. It is therefore not advisable to test and qualify materials for a given application on a natural ageing basis. Accelerated ageing tests are then performed in laboratories in order to elaborate a relatively simple kinetic model describing the material behavior changes and predicting the lifetime from a conventional end-of-life criterion in a reasonable time scale. However, the pertinences of the choice of accelerated ageing conditions, the mathematical form of kinetic model and the end-of-life criterion are rarely discussed and never demonstrated. The use of such an empirical method and, thus, the reliability of its predictions by extrapolation, are highly questionable. In contrast, one can define a nonempirical method for lifetime prediction which would involve at least four elementary steps: (a) identification of elementary ageing mechanisms, and their eventual couplings, through the analytical observations reported in the literature (b) elaboration of a general kinetic model from the previously determined elementary mechanisms (c) determination of model parameters from accelerated ageing tests (d) extrapolation of the considered use properties to (soft) use conditions and lifetime prediction from a critical structural state, using polymer physics. All these steps are very important. Ignoring just one of them can have dramatic consequences on lifetime prediction by extrapolation. The nonempirical method has been applied with success in various industrial applications, for instance, quite recently, in the case of sulphur vulcanized polyisoprene (SVPIP) seals for aerospace applications [1] or carbon black filled polyethylene (CBPE) pipes for the transport of drinking water [2]. The present chapter is devoted to the first step of the nonempirical method. The main polymer (physical and chemical) ageing types, susceptible to take place in practice, will be presented. Particular attention will be paid to embrittlement mechanisms and their corresponding kinetic laws, in order to give to the reader the necessary theoretical bases for a future kinetic modeling.

21.2 Physical Ageing Physical ageing can result from the spatial reorganization of macromolecules (relaxation of enthalpy, volume, orientation or stress, crystallization, etc.), transport phenomena (penetration of a solvent, migration of an additive) and superficial phenomena (cracking in a tensioactive medium).

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Ageing Induced by Structural Reorganization

Liquid-to-glass transition and crystallization are both responsible for polymer solidification in the end of a processing operation. Since they are kinetic phenomena, they lead to an out-of-equilibrium thermodynamic state:

r r

glassy polymers present an excess of unstable conformations and free volume semicrystalline polymers are not totally crystallized, their melting point being largely lower (usually some dozens of degrees) than the equilibrium value.

If, in their use conditions, polymers of either type are subjected to a residual molecular mobility (β motions in glassy polymers, α motions in the rubbery amorphous phase of semicrystalline polymers), they undergo molecular reorganization towards the thermodynamic equilibrium. The main consequences of this ageing type, inappropriately called ‘physical ageing’ since there are other types of physical ageing, best defined as structural relaxation are:

r r r r

a compaction of macromolecules and a loss in enthalpy an increase in yield stress a decrease in creep compliance the variation against time of the previously quoted quantities is auto-retarded, but can proceed for historic durations.

The variation against time of creep compliance J(t) has been the subject of many literature works, in particular in the case of organic glasses [3]. As an example, the general shape of creep curves of samples aged during three different durations: ta , 10 × ta and 100 × ta , is presented in Figure 21.1. The time-shift in creep curves, defined as ac = Log(t1c / t2c ), corresponding to a variation in ageing time aa = Log(t1a / t2a ) is such that aa / ac ≈ 1. In other words, one decade increase in ageing time leads to about one decade increase in the creep characteristic time. It is a general trend of the physical ageing by structural relaxation observed for polymers or molecular organic materials (glucose), as well as for granular solids like sand, or emulsions.

Figure 21.1

Creep curves (in a log-log scale) of samples aged during ta , 10 × ta and 100 × ta in their glassy state.

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Figure 21.2 Shape of stress (σ ) vs. strain (ε) curves before (v) and after (a) structural relaxation. Vertical arrows symbolize sample failure.

The decrease in creep compliance (i.e. increase in Young’s modulus), as well as the increase in yield stress, can be viewed as advantages for many applications; however, these changes are counterbalanced by a catastrophic decrease in ductility / toughness, as schematized in Figure 21.2. The consequence of structural relaxation can be understood as an increase in the fragility of the material. The increase in yield stress (typically, up to 30% of its initial value for a polycarbonate (PC)), the loss in ductility without any variation of the molecular weight distribution, plus the appearance of an endothermic peak close to the glass transition temperature (Tg ) (Figure 21.3) in differential scanning calorimetry (DSC) thermograms allows to unambiguously identify the structural relaxation from any another type of physical ageing in the case of amorphous glassy polymers. Perhaps one of the most dramatic changes in physical properties of amorphous polymers induced by physical ageing is that associated with gas transport properties [4] of polymer membranes with a diffusion coefficient decreasing by nearly one decade. This is a direct effect of the decrease in free volume induced by structural relaxation. CP exo

(a)

(v)

Tg

T

Figure 21.3 Shape of DSC thermograms around the glass transition temperature (Tg ) before (v) and after (a) structural relaxation.

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ε σ1 σ2 σ3 εC

0

σ4 σ5 σ6 t

Figure 21.4 Shape of creep curves (time variation of strain) for different stress levels (σ 1 > σ 2 > σ 3 . . . etc) in the presence of a solvent. Black points indicate the appearance of damage. Their envelope follows a horizontal asymptote for ε = ε c .

21.2.2

Ageing Induced by Solvent Absorption

Solvents plasticize polymers and thus, lead to a decrease in Tg and yield stress (if the polymer is initially ductile). However, the most significant effects, in practice, are observed when the polymer is subjected to a mechanical loading. In this case, plasticization favors damage, in particular by crazing. As an example, let us consider a creep test during which damage is detected by an optical technique. The resulting behavior can be schematized in Figure 21.4. One can see that there is a critical strain εC below which the material does not damage. The value of εC is a function of the polymer−solvent couple. In the simplest cases (moderately polar polymers such as poly(2,6-dimethyl oxyphenylene) (PPO) [5]), εC depends on the parameter of solvent solubility as shown in Figure 21.5. The dashed zone corresponds to values of εC determined in air. In the case of polar polymers

Figure 21.5 Shape of the variation, against parameter of solvent solubility (δ S ), of critical strain (εC ) for PPO [5]. Minimum of curve corresponds to: δ S = δ P , δ P being the parameter of polymer solubility. (◦) crazes; (•) open cracks.

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(for instance, poly(methyl methacrylate) (PMMA)), the behavior can be more complex: the curve εC = f (δ S ) can exhibit several minima. Some vapors, like water or carbon dioxide, play an important role in solvent-induced ageing. Moreover, plasticizers can migrate from one polymer to another and thus generate damage under mechanical loading. In the case of relatively complex pieces, in the presence of vapors, localized damage can be enhanced by the presence of residual stresses due to the processing conditions. The penetration of solvents into a polymer leads to swelling and can also lead to stress gradients induced by hindered swelling during the transient regime of diffusion. Problems of organic matrix composite damage induced by water diffusion have been reported in the literature over the half past century, in particular in aeronautics [6, 7, 8]. 21.2.3

Ageing Induced by Additive Migration

Most technical polymers contain additives like antioxidants, plasticizers or processing agents that do not form chemical bonds with macromolecules. These molecules can migrate within the polymer at a speed depending on the size of the molecules, their solubility and diffusivity within the host polymer among other parameters. Plasticizers (of poly(vinyl chloride) (PVC), for instance) are well-known for their propensity to migrate into any types of polymers. The consequence is an embrittlement of the material. We present first the limiting steps in additive migration. A polymer−additive mixture is out of thermodynamic equilibrium, since the additive concentration in the environment being equal to zero, there is no equality between its chemical potentials in the environment and in the polymer. Additive molecules tend to migrate outside the polymer in order to establish this equality and thus, to reach an equilibrium. This migration is composed of two steps (Figure 21.6): 1 The first step corresponds to the passing of some additive molecules into a medium of molecules close to the polymer surface, that is to say the crossing of the polymer-external medium interface. In a gaseous medium (for instance, atmosphere), the additive evaporation controls the kinetics of this step. In a liquid medium, the additive dissolution within the liquid plays the same role. The external medium can be a polymer in contact with the first one, and again it is the additive dissolution which is the controlling step. A consequence of the last example is that polymers tend to exchange their antioxidants. 2 The additive molecule exchange between the polymer and the external medium leads to a concentration gradient in a region adjacent to the surface of the material. This latter is the ‘driving force’ for the

Figure 21.6 thickness.

Schematization of a two-step-migration of molecular species outside the polymer. L is the sample

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diffusion of additive molecules from the core towards the sample surface. If diffusion obeys the second Fick’s law, the diffusion distance () is proportional to the square-root of time (td ), i.e.: td =

2 D

(21.1)

where D is the diffusion coefficient of additive into the polymer. Note that we treat here Fickian diffusion processes, which may hold essentially far from Tg transition of polymer networks. There are many exceptions to that process and we invite the reader to refer to extensive reviews available in the literature for the treatment of transport of small molecules into polymers [9, 10, 11]. One can see from Eq. (21.1) that td increases rapidly with the sample thickness. From processes (1) and (2) described above, one can distinguish two distinct kinetic regimes of diffusion, depending on which process is the rate limiting step. Regime 1: Evaporation (or dissolution) controlled kinetics In the case of thin samples (fibers, films, coatings, etc.) and high additive diffusivity, additive evaporation is the slowest step and thus controls the global migration kinetics. Its concentration C (into the polymer) decreases proportionally with time (see Figure 21.7, left): dC = −H dt

(21.2)

Figure 21.7 Additive migration governed by evaporation (left) and diffusion (right). Top: shape of weight changes as a function of time. Sample weight (m) corresponds to polymer weight (mp ) plus additive weight (ma ); Bottom: distribution, in the sample thickness, of additive concentration (C) for different times of exposure: t0 < t1 < t2 < t3 . The initial concentration is C = C0 (at t = t0 ).

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The evaporation rate, H, is a decreasing function of the additive molar mass and cohesive energy density. Regime 2: Diffusion controlled kinetics In the case of relatively thick samples (typically few millimetres thick) and low additive diffusivity, diffusion in the bulk controls the global migration kinetics. In the simplest case, the second Fick’s law can be successfully applied. It is found that, in the early period of exposure, the sample weight m decreases proportionally with the square root of time (see Figure 21.7, right): ∂ 2m ∂m =D 2 ∂t ∂z

(21.3)

where z is the spatial coordinate in the sample thickness. Thus, a gradient of additive concentration appears in the sample thickness. Figure 21.7 shows the shape of the time variation of sample mass and concentration profile in the case of evaporation controlled and diffusion controlled additive migration. If the additive molar mass is relatively high, its evaporation (or, more generally speaking, its crossing of the polymer−medium interface) is slow. Then, its concentration at the sample surface takes an intermediate value between C0 and zero. It is necessary to take into account this variation in the boundary conditions, for the resolution of Eq. (21.3). If the additive concentration is high (as in the case of plasticizers), it modifies the polymer properties. Then, its diffusivity becomes concentration dependent. Complications appear when a phase transition takes place during the additive migration. As an example, in the case of diffusion of plasticizers in PVC, a dramatic increase of Tg occurs in regions of low plasticizer concentration. Hence, the Tg profile can have the shape illustrated in Figure 21.8. Naturally, the increase in Tg at the sample surface will lead to a decrease of the

Figure 21.8 Shape of distribution, in the sample thickness, of plasticizer concentration (C expressed in weight fraction) and resulting local glass transition temperature (Tg ) or a plasticized PVC aged at room temperature [12].

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Figure 21.9 Correction of the hypothetic curve (right) presented in Figure 21.7. Real shape of the distribution, in the sample thickness, of plasticizer concentration (C) for different times of exposure: t0 < t1 < t2 < t3 < t4 .

diffusivity in that region, along with, possibly, a change in transport process. The material surface becomes brittle with a ‘soft’ bulk region. One can see that, in the bulk of the sample, the polymer remains in a rubbery state. On the contrary, in the superficial layer of thickness a , the polymer is in a glassy state. Since the coefficient of plasticizer diffusion varies of at least one order of magnitude on both sides of the glass transition temperature, the real gradient will display rather the shape of Figure 21.9. In such cases, the diffusion ‘front’ is very abrupt and the sample weight decreases proportionally with time. On the mechanical point of view, the loss in additives leads to a loss in the specific properties obtained by the introduction of these latter: long-term durability in the case of antioxidants, flexibility in the case of plasticizers, etc. If the additive concentration is high (as with plasticizers), their loss induces a volume shrinkage (of the order of the weight loss). This shrinkage can generate local stresses and thus lead to cracking.

21.3 Chemical Ageing 21.3.1 21.3.1.1

General Aspects The Two Main Families of Chemical Ageing

It is important to distinguish between two types of chemical ageing processes: on the one hand, those leading to change in lateral groups of macromolecules, without affecting the macromolecular skeleton; on the other, those leading also to a change in the macromolecular skeleton. Indeed, in the first case no significant change in the mechanical properties at low conversion ratio of the reaction is produced, whereas the second type induces an important change of mechanical properties, even at low conversion ratio. Mechanisms Leading to a Change in Lateral Groups Halogenation and sulfonation reactions are susceptible to substitute hydrogen atoms by various groups, but such mechanisms are rarely encountered in use conditions. At the opposite, oxidation reactions are very common. These reactions may affect also the

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macromolecular skeleton but, at this stage, only their consequences on the macromolecule chemical structure will be considered. Examples of lateral groups produced by oxidation are the following: OOH

O2 CH2

OH or

CH

CH

O or

C

A hydrocarbon group, moderately polar, is transformed into a polar (ketone) or highly polar (alcohol, acid) oxygenated group. If this change occurs in a polymer initially rich in polar groups, it has limited consequences on physical properties. However, if this change occurs in a polymer initially not very polar (such as polyolefins or hydrocarbon elastomers), it induces a significant variation of all the physical properties depending on the macromolecule chemical structure, in particular: coloring, dielectric permittivity, refractive index, surface energy and wettability. Some mechanisms lead to the formation of colored species. As an example, in the case of PVC, one of them is the sequential (zip) elimination of hydrogen chloride: S

CH2 S

CHCl CH

CH2 CH

CHCl CH

CH

CH2 CH

CHCl CH

+

n HCl

where S is a structural irregularity (for instance, a tertiary or an allylic chloride formed during polymer synthesis) destabilizing the neighboring monomer unit. Sequential elimination is favored because allylic chlorides are especially unstable. It leads to the formation of conjugated polyenes: CH

CH

j

with j ranging typically from 1 to 20. The absorption spectra of polyenes in the UV-visible range are schematised in Figure 21.10 [13]. As a general rule, for j ≥ 6, polyenes absorb in the UV-visible range, and the absorption spectrum is shifted towards long wavelength values as the number of conjugated double bonds increases. The polymer appears yellow for j ≤ 8, violet−purple for j = 12−15, and brown to black for higher values of j. Since the resulting chemical species present a very high coefficient of molar absorptivity, of the order of 105 l.mol−1 .cm−1 , coloration appears at low, and sometimes at very low, conversion ratio of the reaction. In the case of aromatic polymers, oxidation leads to chromophoric species, in particular quinoleic species absorbing in the near UV and violet range: O O2

O

Thus, these polymers (polystyrene (PS), polycarbonate (PC), poly(ethylene terephthalate) (PET), styrene crosslinked polyesters (UP), epoxy, etc) undergo yellowing during exposure to oxidation conditions. Oxidation of saturated hydrocarbon polymers (for instance, polyethylene (PE) and polypropylene (PP)) can also lead to colored species if the material contains initially traces of metallic impurities susceptible to form colored complexes with some oxidation products such as carbonyl groups.

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Figure 21.10 Shape of UV-visible spectra for conjugated polyenes as a function of the conjugation degree j indicated in the figure [13]. The absorption wavelength peak (λ) and molar absorptivity (ε) increase with conjugation length.

Finally, some stabilisation mechanisms, in particular with aromatic amines or hindered phenols, are susceptible to leading also to highly-colored species (due to antioxidant reaction). However, these latter can be easily distinguished from colored species belonging to macromolecules, since they are extractible by using an adequate solvent. Mechanisms Leading to a Change of Macromolecular Skeleton One can distinguish two distinct types of ageing induced at the level of the macromolecular skeleton: chain scission and crosslinking (Figure 21.11). Of course, both types can take place simultaneously. If S and X are respectively the numbers of moles of chain scission and chain welding per mass unit, one can write [14]: 1 1 − =S−X Mn Mn0 1 S 1 − = − 2X MW MW 0 2

(21.4) (21.5)

- Random chain scission:

- Chain welding (crosslinking):

Figure 21.11

Schematization of the two types of changes in polymeric chain skeleton during chemical ageing.

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where M n0 , M W0 , Mn and MW are respectively the number and weight average molar masses before and after ageing. One can see that, if S > 4X, chain scission predominates. Alternatively, if S < 4X, crosslinking predominates. As long as the material remains 100% soluble, molar mass measurements (by steric exclusion chromatography (CES), rheometry, etc.) can be carried out to evaluate such chemical changes. Saito’s theory (Eqs (21.4) and (21.55)) allows the determination of S and X. However, if crosslinking predominates, the material reaches a gel point (MW → ∞ and X → Xg , whereas Mn remains finite) and an insoluble fraction appears. Beyond the gel point, the determination of soluble fraction (ws ) or elastic properties in rubbery state are the most common analytical ways of characterization. Then, Charlesby’s theory can be used to determine S and X: 1/2

wS + wS

=

S 1 − 2X MW 0 X

(21.6)

This relationship is easily applicable if S and X are linear functions of time, as is generally the case in radiochemical ageing at relatively high dose rates. Consequences of a Random Chain Scission and (21.5) lead to:

In linear polymers, in the absence of crosslinking, Eqs (21.4)

MW (PI 0 − 2) MW 0 PI 0 − 2 i.e.: PI = 2 + 1 + SM W 0 /2 PI = 2 +

(21.7)

where PI 0 and PI are the polydispersity indexes respectively before and after ageing. Since MW decreases during degradation, IP tends towards 2 at long ageing time (large conversion ratio), see Figure 21.12. Thus, if PI is initially close to 2, it will remain constant during degradation by random chain scission: PI 0 ≈ PI ≈ 2. Such a simplification is current in the field of polymer chemical ageing.

Figure 21.12 Shape of the variation of polydispersity index (PI) as a function of ageing time for a linear polymer undergoing a random chain scission: (a) PI0 > 2. (b) PI0 < 2.

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Figure 21.13 Shape of the variation of toughness (G1C ) for a linear polymer as a function of its weight average molar mass (MW ).

Fracture properties of virgin linear polymers, for instance their toughness G1C or their ultimate elongation εR , increase abruptly of several decades when their molar mass MW becomes higher than a critical molar mass MF (cf. Figure 21.13), this latter being closely related to the rheological molar mass MC [15]. One can thus expect that, during degradation by random chain scission, the fracture properties of an initially ductile/tough linear polymer decrease catastrophically when its molar mass becomes lower than the MF value. Below MF , the polymer behaves like a wax or an eggshell, depending on its stiffness. It can no longer be used in mechanical applications owing to its extremely high brittleness. Now, let’s consider a random chain scission process occurring in a linear polymer with a constant rate rS . The number of chain scission is just: S = rS t

(21.8)

Assuming that PI 0 ≈ PI ≈ 2, Eq. (21.4) can be rewritten as: MW =

2MW 0 2 + r S MW 0 t

(21.9)

The variations in time of the molar mass and ultimate elongation are schematized in Figure 21.14. These results call for the following comments: 1 Whereas S is a linear function of time, MW is an hyperbolic function of time; 2 MW is a continuously decreasing function of time, whereas εR falls abruptly when MW = MF . As a general rule, embrittlement occurs prematurely. As an example, let’s consider the two application cases that are hydrolysis at ambient temperature of PA11 pieces in offshore applications and radio-oxidation at ambient temperature of PE used as electrical insulation in cables for nuclear power plants. For such polymers, typical molar masses (M W0 and MF ) are reported in Table 21.1. From these values, the critical number of chain scissions per mass unit, responsible for the polymer embrittlement (εR / εR0 < 50%) has been calculated

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Figure 21.14 Shape of the variation in ageing time of weight average molar mass (MW ) and ultimate elongation (εR ) for a linear polymer subjected to a random chain scission process.

from: SF =

2 2 − MF MW 0

(21.10)

It is compared to the concentration of reactive sites susceptible to undergo a chain scission process (i.e. hydrolyzable amide groups for PA11 and oxidizable methylene groups for PE) in each polymer: Cm =

1 m

(21.11)

where m is the molar mass of the monomer units: m = 0.183 kg.mol−1 for PA 11, and 0.014 kg.mol−1 for PE. The resulting conversion ratio YS (at the molecular scale) of the chain scission process has been also reported in Table 21.1. YS =

Cm SF

(21.12)

Table 21.1 Embrittlement criteria for some common polymers. MW0 and MF are respectively initial and critical weight average molar masses; SF is the critical number of chain scission; Cm is the concentration of potentially reactive sites and YS is the resulting conversion ratio (at the molecular scale) of the chain scission process. Polymer

Chemical ageing

PA11 PE

Hydrolysis Oxidation

MW0 (kg.mol−1 )

MF (kg.mol−1 )

SF (mol.kg−1 )

Cm (mol.kg−1 )

YS (%)

50 150

30 70

2.7 10−2 1.5 10−2

5.5 71.4

0.5 0.02

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One can see that embrittlement occurs for a very low number of chain scissions and thus, a very low conversion ratio of reaction usually inaccessible by conventional spectroscopic techniques (infrared spectroscopy, nuclear magnetic resonance spectroscopy, etc.). This means that molar mass characterization (by steric exclusion chromatography or rheometry) is much more sensitive than such techniques for characterizing chain scission in the hypothesis that it is the only process at play. Elastic and yield properties are almost not altered at such low conversion ratio. From then on, one can understand that chemical ageing-induced embrittlement constitutes a real domain of interest for many researchers. Whatever the importance of other use properties, no application can tolerate that the geometrical integrity of pieces be altered because of cracking. Photo-oxidation affects only a thin superficial layer directly exposed to solar radiation, for instance a few dozens of micrometers in the case of epoxies [16]. In such chemical ageing types, local embrittlement can generate a network of microcracks without altering the global fracture behaviour of pieces. However, micro-cracking alters the surface aspect of pieces and may therefore have a significant economical impact for some applications, for instance in the case of varnishes of automotive bodies [17], for which aesthetic questions are particularly important. In polymer networks (vulcanized elastomers, thermosets), random chain scission leads to a destruction of elastically active chains (EAC). Typically, at low conversion ratios of reaction, for a quasi-ideal network with a functionality f ≥ 4: ν = ν0 − S

(21.13)

where ν 0 and ν are the concentrations in EAC respectively before and after ageing. These quantities can be determined from measurements of elastic properties in rubbery state:

ν=

E G = 3ρRT ρRT

(21.14)

where E and G are respectively the Young’s and shear moduli (expressed in Pa), ρ is the volumic mass (in kg.m−3 ), R the perfect gas constant and T the absolute temperature (in K).

S=

1 1 (E 0 − E) = (G 0 − G) 3ρRT ρRT

(21.15)

Chain scission leads to a decrease of elastomers modulus. On the contrary, they have no direct consequences on glassy thermosets modulus. Fracture properties tend to decrease catastrophically, but structure−property relationships are not yet well known in this domain. Consequences of Crosslinking In linear polymers, in the absence of chain scission, MW and Mn increase, PI increases, and PI and MW → ∞ when the material reaches a gel point (X → Xg ):

Xg =

1 2MW 0

(21.16)

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One can see here also that gelation corresponds to a very low conversion ratio of reaction. Since it is necessary that two reactive sites encounter to weld two macromolecules, the conversion ratio at gel point is: Yg =

1 MW 0

(21.17)

As an example, in the case of a PE having an initial molar mass M W0 = 150 kg.mol−1 , one obtains: Yg ≈ 6.6 10−3 mol.kg−1 , whereas the concentration in potentially reactive sites (methylene groups) is 104 times higher. Of course, some properties, in particular the flow properties in melt state, vary far before the gel point. In particular, the viscosity in melt state increases rapidly and the Newtonian plateau tends to disappear. The ultimate elongation decreases progressively, whereas the stress at break increases, at least in the early period of exposure, for elastomers. In the case of thermosets, glass transition temperature and thus, yield stress increase, making yielding less and less competitive with brittle fracture. 21.3.1.2

Diffusion Phenomena in Chemical Ageing

Oxidation and hydrolysis, both being very common chemical ageing types, involve the penetration of a small reactive molecule (oxygen, water) initially present in the surrounding environment, into the material. One can thus define two characteristic times: one for diffusion, other for chemical reaction, for elaborating a simplified approach of the problem. In the case of diffusion, the characteristic time can be defined such as: tD =

L2 D

(21.18)

where L is the sample thickness and D is the coefficient of reactive molecule diffusion. In the case of reaction, one can use the following definition: tR =

C rC

(21.19)

where C is the equilibrium value of reactant concentration into the polymer, and rC is the rate of reactant consumption by the chemical reaction in the superficial layer of the material where its concentration is actually C. From the equality of the critical times for diffusion and reaction, one can deduce a critical distance LC in the material as:  Then: L C =

DC rC

1/2 (21.20)

If the sample thickness is L  LC , diffusion is rapid enough against reaction and, at any time t, the whole sample thickness is saturated in reactant. rC is depth independent. On the contrary, if the sample thickness is L LC , all the reactant is consumed in a superficial layer of thickness LC and the material core remains intact. The comparison of the global degradation rates of samples of different thicknesses allows the construction of the graph presented in Figure 21.15, which allows the determination of the value of LC . One can distinguish two distinct kinetic regimes of reaction, depending on the material thickness, as depicted in Figure 21.15:

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Figure 21.15 Shape of the variation, against reciprocal sample thickness (L−1 ), of global degradation rate (r) in the hypothesis that both faces of the sample are exposed to diffusing species.

1 Regime (I), corresponding to thin samples, is not controlled by diffusion. The global degradation rate is thickness independent. 2 Regime (II) is diffusion controlled. The global degradation rate is proportional to the reciprocal thickness since the thickness of the degraded layer LC is thickness independent. The value of the boundary between the two regimes is 2LC if the penetration of the reactive molecules takes place through both sample faces, as is usually the case. For a more rigorous and more detailed approach to the problem, it is necessary to derive the relationship rC = r(C), expressing the rate of reactant consumption as a function of its concentration in relatively thin samples (L  LC ), from a kinetic analysis of the corresponding mechanistic scheme. The rate of variation of the local reactant concentration in a relatively thick sample (L 2LC ) can be determined from the following balance equation: ∂ 2C ∂C = D 2 − r (C) ∂t ∂z

(21.21)

where z is the spatial coordinate for an infinite plate (e.g. the depth beneath the sample surface). The numerical resolution of this equation gives access to the distribution, in the sample thickness, of reactant concentration and its variation against time: C = f (z, t). In the case of hydrolysis, r(C) takes, in general, the mathematical form of a pseudo-first-order reaction: r (C) = KC

(21.22)

On the contrary, in the case of oxidation processes, r(C) can be expressed, in a first approximation, by an hyperbolic function: r (C) =

aC 1 + bC

(21.23)

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where a and b are parameters to be expressed as functions of the rate constants of elementary reactions constituting the radical chain oxidation reaction. As a general rule, chemical reactions are temperature dependent. In the scale equation (Eq. (21.21)) issued from a simplified approach of the problem, the equilibrium concentration of reactant C is generally temperature independent, whereas the diffusion coefficient D obeys an Arrhenius law:   ED D = D0 exp − RT

(21.24)

where D0 and ED are respectively the pre-exponential factor and the activation energy for reactant diffusion. It has been shown that the temperature variation of the rate of reactant consumption in the superficial layer of the material, rC , can be satisfyingly approximated by an Arrhenius law:   Er rC = rC0 exp − RT

(21.25)

where rC0 and Er are respectively the pre-exponential factor and the activation energy for reactant consumption. Finally, one can see that LC obeys also an Arrhenius law:   EL L C = L C0 exp − RT 1/2  D0 C with: L C0 = and rC0

(21.26) EL =

1 (E D − Er ) 2

In most of the cases of thermal oxidation and hydrolysis, ED < Er so that EL is negative. Thus, LC is a decreasing function of temperature. In the case of irradiation-induced chemical ageing, for instance radio and photo-oxidations, C and D are light intensity independent and Eq. (21.24) remains valid. At the opposite, it can be demonstrated that the variation against light intensity I of rC can be satisfyingly approximated by a simple power law: rC ∝ I α

(21.27)

with α = 1/2 in the simplest models. Thus, the thickness of the oxidized layer is a slowly decreasing function of dose rate given by: L C ∝ I −α/2

(21.28)

The fact that the thickness of degraded layer is, in general, a decreasing function of the ‘severity’ of ageing conditions has been systematically observed by many authors. 21.3.1.3

Consequences of Superficial Degradation on Use Properties

Degradation gradients, resulting from a diffusion control of chemical reaction kinetics, play an important role from a mechanical point of view. Schematically, an initially ductile / tough and homogeneous polymer sample is progressively transformed into a binary structure, the sample core remaining ductile / tough, whereas the superficial layer becomes brittle and thus, highly sensitive to cracking. A superficial crack can cross rapidly

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Figure 21.16 Schematization of chemical ageing-induced cracking. Sample zones where the polymer chemical structure has been changed are represented in grey. (0) Initial (virgin) sample; (I) Superficial degraded layer below its embrittlement threshold; (II) Superficial degraded layer close to its embrittlement threshold; (a-III) Crack having reached the skin−core interface and having blunted; (b-III) Crack having crossed the interface; (a-IV) and (b-IV) Propagation of oxidation front towards deeper layers.

the whole superficial layer thickness and reach the skin−core interface. At this stage, two scenarii can take place (Figure 21.16): (a) crack tip blunts and remains restricted to the sample superficial layers (b) crack crosses the interface and propagates in the core. Since cracks constitute a preferred path for small reactive molecules penetration, the degradation front moves towards deeper layers. A secondary degraded layer, of the same thickness as the primary one, forms beyond the crack tip, etc. Following this scheme, failure will ultimately occur, even without any external loading. In a first approach, one can consider that the brittle superficial layer is equivalent to a ‘natural’ notch with the same depth. Fracture mechanics [18] predicts that there is a critical notch depth below which the notch does not initiate the material failure. This depth depends on the material toughness, this latter depending, in turn, on the rate of crack propagation. As an example, in the case of PE oxidation, this depth is of the order of magnitude of 100 μm [19]. It is thus expected that rapid chemical ageing, leading to very small thicknesses of degraded layer, will have less effect on the material fracture behavior than slower ones. This general trend has been also observed by many authors. 21.3.2

Mechanistic Schemes

Chemical ageing is a complex phenomenon, but decomposable into a more or less high number of elementary reactions obeying relatively simple kinetic laws: Arrhenius law, proportionality to light intensity, etc. Although chemical ageing proceeds through molecular scale (monomer unit) processes, we have shown in the previous sections that use properties of materials depend also on structural events occurring at the upper scales, in particular at the macromolecular (macromolecular skeleton) and the macroscopic ones (skin−core structure). The objective of the present section is to give a rough description of the main chemical reactions occurring in a polymer during exposure to an aggressive chemical environment.

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Table 21.2 Orders of magnitude of the dissociation energy (ED ) of main polymer chemical bonds.

21.3.2.1

Chemical bond

ED (kJ.mol−1 )

aromatic C− −C C− −F aromatic C− −H aliphatic C−−H aliphatic C−−C C− −O C− −Cl C− −Si C− −N C− −S S− −S O−−O

510 470 465 325−425 300−380 340 320 300 290 275 260 150

Intrinsic Chemical Stability

In the absence of reactive molecules (water, oxygen) in the environment of exposure, polymers can decompose because of the own material instability, this latter being governed by thermochemical factors. It involves, in general, radical reactions initiated by the breakdown of the weakest chemical bonds (i.e. characterized by the lowest dissociation energy) of macromolecules. Aromatic C−−C and C−−H bonds and C−−F bonds (dissociative energy ED > 450 kJ.mol−1 ) are the most stable (Table 21.2). Polymers only composed of such chemical bonds are called, on purpose, ‘thermostable polymers’. Polyimides and fluorinated polymers, e.g. polytetrafuoroethylene (PTFE), are examples of thermostable materials. Aliphatic C−−C bonds are all the more unstable than the carbon atoms are more substituted, see for instance two series of bonds ranked by decreasing dissociation energy: CH3 CH3

CH3

CH2

>

CH3 CH2

>

C

CH3

>

CH3

CH3

CH

C

CH3

CH3

CH2

CH2

>

CH2

CH

>

CH3 CH3

>

C

C

CH3 CH3

Thus, polymers containing tetrasubstituted carbons, such as polyisobutylene (PIB), poly(methyl methacrylate) (PMMA), poly(α methyl styrene) (PαMS), etc, are particularly unstable (ED = 300−320 kJ.mol−1 ). Also, polymers containing heteroatoms, for instance poly(oxy methylene) (POM), poly(ethylene terephthalate) (PET), poly(vinyl chloride) (PVC), sulphur vulcanized polyisoprene (SVPIP), etc, are also particularly unstable (Table 21.2). However, the global thermal stability does not depend only on the presence of weak chemical bonds. It depends also on the ability of radicals to initiate a chain propagation reaction. As an example, in the case of polyethylene (PE), radicals deactivate easily by disproportionation: CH2

CH

+

CH2

CH

CH2

CH2

+

CH

CH

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In the case of polymers containing tetrasubstituted carbons, as well as POM, radicals initiate a sequential (zip) elimination of monomer units: M

M

M

M

M

M

M

M

M

M

M

+

+ M

M

M +

M

M

M

M

This reaction is called a ‘depolymerization’. This is actually a weak point as polymers which depolymerize have the lowest thermal stability. PVC undergoes a sequential (zip) elimination of its lateral groups: S

CH2

CHCl

CHCl

CH2

S

CH

CH

CH2

S

CH

CH

CH

CHCl

CH

+ HCl

+ HCl

etc ...

where S is a structural irregularity (for instance, a tertiary or an allylic chloride formed during polymer synthesis) making the first elimination step easier. In the case of step-by-step processes, such as PE thermal degradation, there is no well-known way of stabilization by acting on the macromolecular structure. On the contrary, in the case of chain reactions (depolymerization, sequential elimination of volatile products like, for instance, HCl, etc.), stabilization can be envisaged through several routes: 1 Suppressing the ‘weak points’ by chemical modification during polymer synthesis (blocking of terminal groups in the case of POM) or by stabilization reaction during the processing operation (using stabilizers for PVC). 2 By interrupting efficiency the propagation by using an adequate co-monomer, for instance in the case of POM: CH2

O

CH2

O

CH2

O

CH2

CH2

O

CH2

O

CH2

O

The insertion of ethylene units (as underlined in the above scheme) in the macromolecule by copolymerization allows slowing down depolymerization. POM copolymers can be easily distinguished from POM homopolymer: the former presents a melting point at about 165 ◦ C, i.e. about 10−20 ◦ C lower than the latter. 21.3.2.2

Oxidative Ageing

General Aspects In most of the cases of thermal, radiochemical or photochemical ageing, the main ageing cause is the polymer oxidation by atmospheric oxygen. In all cases, degradation propagates through a radical chain reaction, the only difference being the initiation step. Thus, polymer oxidation kinetics can be

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described satisfactorily by a general mechanistic scheme called ‘standard scheme’, established in the 1940s by the RAPRA co-workers [20]: (1) Polymer→ μ P◦ (2) P◦ + O2 → PO2 ◦ (3) PO2 ◦ + Polymer → Per + P◦ (4) P◦ + P◦ → inactive products (5) P◦ + PO2 ◦ → inactive products (6) PO2 ◦ + PO2 ◦ → inactive products + O2

(r1 ) (k2 ) (k3 ) (k4 ) (k5 ) (k6 )

In the ‘standard scheme’ of oxidative ageing, (1) corresponds to the initiation step, (2) and (3) to propagation steps; (4)–(6) to termination steps. P◦ and PO2 ◦ refer respectively to alkyl and peroxy radicals; Per refers to peroxide groups; μ is the formation yield of radicals in the initiation step; r1 is initiation rate and ki are rate constants. The initiation step (1) can be thermal, photochemical or radiochemical. It can result from a reaction involving directly lateral groups of monomer units (radiochemical ageing), or unstable pre-existing products like hydroperoxides and peroxides Per (particularly in thermal and photochemical ageing). The propagation reaction is composed of two elementary steps (2) and (3): (2) Oxygen addition to alkyl radicals: It is a very fast process and thus, practically structure- and temperatureindependent. The corresponding rate constant is very high: k2 = 108 –109 l.mol−1 .s−1 [21]. (3) Peroxyl radicals reaction with the polymer: this is a largely much slower process, which is structuredependent. In saturated hydrocarbon polymers, for instance polyethylene (PE) and polypropylene (PP), it corresponds to hydrogen atom abstraction. In this case, Per is a hydroperoxide group (POOH). The corresponding rate constant is very low: k3 = 10−3 –10−1 l.mol−1 .s−1 at ambient temperature (cf. Table 21.3). In polyenic elastomers, for instance polybutadiene (PBD) and polyisoprene (PIP), step (3) can be an addition to double bonds. Thus, Per can be a peroxide group POOP. The corresponding rate constant is also very low: for instance k3 = 10−1 –10 s−1 at ambient temperature in the case of the intramolecular addition (cf. Table 21.3).

Table 21.3 Propagation reactions of oxidation in some common hydrocarbon polymers and corresponding rate constant (k3 ) value at ambient temperature. PP: polypropylene; PE: polyethylene; PDB: polybutadiene; PIP: polyisoprene. Propagation

Polymer

k3 (l.mol−1 .s−1 or s−1 )

Source

Hydrogen atom abstraction

PP PE PBD PIP PBD PIP

1.0 × 10−3 2.4 × 10−3 4.9 × 10−3 5.2 × 10−2 6.1 × 10−1 2.7

[22] [22] [23] [1] [23] [1]

PBD PIP

5.8 × 10−4 −

[23] [1]

Intramolecular addition to double bonds Intermolecular addition to double bonds

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In the absence of antioxidants, the terminations of radicals are bimolecular processes. At relatively low temperature, like ambient temperature, termination rate constants of reactions given in the above scheme classify in the following order [24]: k4 > k5 k6

(21.29)

(corresponding activation energies classify in the reverse order), mainly because P◦ radicals can propagate by hydrogen abstraction: P◦ + PH → PH + P◦ . This transfer reaction does not influence the whole oxidation kinetics (except for polyenic elastomers), but provides a simple explanation to the relatively high mobility of alkyl radicals and thus, the high k4 value. However, at moderate to high temperature, reactions (5) and (6) lead mainly to unstable peroxides (POOP). As a result, Eq. (21.29) is no longer valid. When T > 200 ◦ C, for instance, it is rather observed [25, 26]: k4 > k5 ≈ k6

(21.30)

Over the past half century, the most delicate and debated subject has been initiation. Initiation processes can be very varied and complex. They depend on the polymer nature (PH) and the way energy is brought to the system (by temperature or radiation). Though intermediate steps are not always totally known (often due to unsuccessful identification of chemical evolution during this step), the common output of this (these) step(s) is the formation of alkyl radicals P◦ . For sake of simplicity, we consider the case of oxygen excess (relatively thin polymer sample exposed to a relatively high oxygen pressure). In this case, all the P◦ radicals are almost instantaneously transformed into PO2 ◦ and then their probability to participate to reactions other than reaction (6), in particular reactions (4) and (5), is negligible. As a result, the ‘standard scheme’ reduces to four reactions (1, 2, 3 and 6). One can distinguish two important initiation processes, as outlined below. Case 1: Initiation at a Constant Rate In the case of radiochemical ageing (high energy provided), the main source of radicals is polymer radiolysis, i.e. the breakdown of lateral bonds of monomer units. In the case of PE, radiolysis leads to the formation of very reactive radicals H◦ which recombine rapidly by hydrogen atoms abstraction:

H

CH2

CH2

CH2

+

CH2

CH2

+ hν CH2

CH2

CH

CH2

+

H

CH2

CH

CH2

+

H2

Thus, the corresponding balance initiation can be written: (1)

PH

→ P◦ + 12 H2

(r 1 )

where the initiation rate r1 is proportional to the dose rate I according to: r1 ≈ 10−7 G 1 I

(21.31)

and G1 is the radical yield expressed in number of radicals P◦ per 100 eV absorbed, of the order of magnitude of 1−10 for saturated hydrocarbon polymers [27].

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Figure 21.17 Shape of oxidation kinetic curves in the case of initiation (a) at constant rate (e.g. radiochemical ageing); (b) by bimolecular Per decomposition (thermal ageing); and (c) by unimolecular Per decomposition (case of many photochemical ageings).

In this case, the reaction reaches rapidly a steady state and oxidation products accumulate with a constant rate (cf. Figure 21.17). One can easily demonstrate that initiation and termination products form with a rate proportional to r1 . As an example, if ketones are formed with a yield γ K in termination: r K = γ K k6 [PO◦2 ]2 =

γK r1 2

(21.32)

On the contrary, propagation products form with a rate proportional to the square root of r1 , for instance: rPer = k3 [PO◦2 ][PH] = k3 [PH]



r1 2k6

1/2 (21.33)

Finally, oxygen is absorbed with a rate: ◦

r O2 = k2 [O2 ][P ] −

k6 [PO◦2 ]2

  r1 r1 1/2 + k3 [PH] = 2 2k6

(21.34)

Case 2: Initiation by Decomposition of Peroxides In the case of thermal and photochemical ageing (lower energy provided), the problem is significantly more complex. Chemical bonds of common industrial polymers have rarely a dissociation energy lower than 260 kJ.mol−1 (Table 17.2) and thus decompose only at high temperatures (typically, at T > 250 ◦ C) or at high irradiation intensities (for instance, under γ irradiation in nuclear environments). However, in the previous sections it has been seen that oxidation leads to the formation of two main propagation products: hydroperoxides POOH and peroxides POOP, noted Per, of which the activation energy of O−O bond is very low: ED ≈ 150 kJ.mol−1 . Such chemical groups are thermally and photochemically unstable, in particular in use conditions.

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It cannot be totally excluded that an ‘extrinsic specie’ generates also primary radicals (e.g. decomposition of structural irregularities or direct oxygen−polymer reaction). But it can be easily demonstrated that its contribution to initiation is very limited: the corresponding initiation rate is initially very low and vanishes rapidly (as soon as the ‘extrinsic specie’ concentration vanishes). As a result, in all cases, initiation by Per decomposition rapidly becomes the main source of radicals. From a kinetic modeling point of view, the following approach is usually adopted: initially present ‘extrinsic specie’ is replaced by a kinetically equivalent initial Per concentration: [Per]0 . In this case, initiation involves two reactions: (1) δPer → αP◦ + βPO◦2

(k 1 )

with: α = 2 and β = 0 for unimolecular decomposition (δ = 1) α = 1 and β = 1 for bimolecular decomposition (δ = 2) Then, the initiation rate depends on Per concentration: r1 = k1 [Per]δ

(21.35)

Such an oxidation mechanism is called a ‘close-loop mechanism’, because it produces its own initiator product (Per). The resulting kinetic curves present an induction period followed by a sharp auto-acceleration, preceding a steady state. The auto-acceleration step is much more progressive in its initial phase when Per decomposition is unimolecular (the case of many photochemical ageings) (cf. Figure 21.17). According to analytical models, the duration of the induction period is given by: 5 (unimolecular decomposition) 2k1 1 − ln Y0 ti ≈ (bimolecular decomposition) K ti ≈

(21.36) (21.37)

 k1 1/2 [Per]0 and Y0 = k6 [Per] S Moreover, in steady state, the Per concentration is: 

where K = k3 [PH]

δ

[Per] = −δk1 [Per] +

k3 [PO◦2 ][PH]



k 2 [PH]2 = 32 δ k1 k6

1/δ (21.38)

and the rate of oxygen consumption is: rO2 = k2 [O2 ][P◦ ] − k6 [PO◦2 ]2 =

2 k32 [PH]2 δ 2 k6

(21.39)

The kinetic analysis of oxidation mechanisms shows that the kinetic behavior depends, in both cases 1 and 2, on two main factors: 1 An extrinsic factor, that is to say an external factor to polymer structure: initiation rate r1 or initiation rate constant k1 ; k3 [PH] . 2 An intrinsic factor characterizing the polymer oxidisability: ratio √ k6

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Table 21.4 Orders of magnitude of the dissociation energy (ED ) of CH bonds. ED (kJ.mol -1) 465

C−H bond

CH3 CH2

414 393 378

CH2

CH O

CH 2

or

C

CH

CH2

N

CH 2

376 335

According to some authors [22], there is a linear relationship between Log k3 and the dissociation energy ED of CH bonds. Thus, in a first approach, polymer stability to oxidation can be roughly estimated from the reactivity of the CH bonds involved (cf. Table 21.4). The following global trends can be deduced: 1 Polymers without CH bonds such as poly(tetra fluoro ethylene) (PTFE), or containing exclusively aromatic CH bonds such as poly(ether ether ketone) (PEEK), poly(ether sulfone) (PES) and polyimides (PI), are stable to oxidation. 2 Polymers containing exclusively methyl CH bonds such as poly(dimethyl siloxane) (PdMS), or containing methyl and methylene CH bonds such as: poly(methyl methacrylate) (PMMA), polycarbonate (PC) and polyethylene (PE), are moderately stable. 3 Polymers containing methyne CH bonds such as polypropylene (PP), or methylene CH bonds in a α position of an heteroatom such as: poly(methylene oxide) (POM), polyamides (PA) and amine crosslinked epoxy (ACE), are relatively unstable. 4 Finally, polymers containing allylic CH bonds such as polyisoprene (PIP) and polybutadiene (PBD), are highly unstable. The above section has treated ageing kinetics only through intrinsic stability of polymeric chains. There are actually other factors of determining importance when dealing with polymers in real conditions, with two examples given: Oxygen diffusivity play an important role through provision of one of the necessary reactants: It is about three orders of magnitude higher in elastomers than in glassy polymers. Thus, at equal reactivity, glassy polymers will appear much more stable because their oxidized layer will be much thinner. At least another factor plays an important role for use properties: the polymer sensitivity to chemical events resulting from oxidation. As an example, mechanical embrittlement occurs in polypropylene (PP) for a number of chain scissions ten times higher than in an amorphous polymer. Thus, at equal reactivity, according to this mechanical endlife criterion, PP will be ten times less stable than an amorphous polymer. Resulting Macromolecular Changes If oxygen is in excess (no termination involving alkyl radicals), in general, chain scission is largely predominant over crosslinking. The β scission of alkoxy radicals PO◦ is probably the most frequent source of chain scissions in polymer oxidation. PO◦ radicals come from the decomposition of Per groups: δPer → αPO◦ + βPO2 ◦

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or from the nonterminating bimolecular combinations of peroxy radicals PO2 ◦ : PO◦2 + PO◦2

→ PO−−O−−O−−OP

[PO◦◦ OP]cage

→ inactive products

[PO◦◦ OP]cage

→ 2PO◦

→ [PO◦◦ OP]cage + O2

β scission can be schematized as follows: R C

R

O2 C

C

C

H

R C

C

C

+

C

C

O

O

It leads to an alkyl radical and a carbonyl group. If R = H (case of polyethylene (PE)), this latter is an aldehyde, whereas if R = CH3 (case of polypropylene (PP)), it is a ketone. Note that this reaction is in competition with hydrogen abstraction: PO◦ + PH

→ P−−OH + P◦

Thus, if the chain scission process occurs with a yield γ S in initiation or non terminating combination, the corresponding chain scission rate is: γ S6 r1 2 γ S k 2 [PH]2 r S = γ S k1 [Per]δ = 2 3 δ k6 r S = γ S6 k6 [PO◦2 ]2 =

(radiochemical ageing)

(21.40)

(thermal and photochemical ageing)

(21.41)

Crosslinking predominates only in some polyene elastomers, for instance in polybutadiene (PBD), in which the addition of peroxy radicals to double bonds is, partly, intermolecular:

PO2° +

CH2

C

CH

CH2

CH2

C

CH

CH2

O O

P

Stabilization Chain reactions, such as radical chain oxidation, especially if they produce their own initiator product, present a serious drawback. It can be called the ‘butterfly effect’: small causes lead to great consequences. However, this drawback is transformed into advantage if one envisages a stabilization by additives, in particular antioxidants. Since oxidation starts very slowly, from a low concentration of reactive species (Per, radicals), and then auto-accelerates, one can envisage to inhibit it with a low concentration of adequately chosen antioxidants. On the contrary, if an extrinsic initiation, with a constant rate, occurs, antioxidants lose their efficiency [28]. In the case of a ‘closed-loop’ oxidation mechanism, there are two possible ways of stabilization: (a) A reduction of initiation rate by decomposing hydroperoxides by a nonradical way: (7) POOH + Dec → inactive products (k7 )

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Figure 21.18 Shape of DSC thermograms in pure oxygen at 190 ◦ C for nonstabilized polyethylene (0), and the same polyethylene stabilized by 0.3% in weight of a phenol−phosphite mixture (S) [32]. In both cases, heating from ambient temperature to 190 ◦ C has been made under nitrogen and oxygen has been admitted in the cell after system thermal equilibrium.

In this case, Dec is an organic sulphur or phosphite. Antioxidants belonging to this first family can be called ‘preventive antioxidants’. (b) A reduction of the propagation rate, i.e. an increase of the termination rate, by capturing radicals (P◦ and PO◦2 ): (8) P◦ + >NO◦ → >NO−P (k8 ) In this case, R–NO◦ is a nitroxy radical. This latter can be directly incorporated to the polymer during processing or formed by oxidation of an hindered amine (HAS) preliminary incorporated. (9) PO◦2 + AH → POOH + inactive products (k9 ) Here, AH is a hindered phenol or an aromatic amine. Such molecules give easily their hydrogen atom to peroxy radicals, because its dissociation energy is smaller than that of hydrogen atoms of saturated hydrocarbon polymers. As an example, in the case of hindered phenols, the dissociation energy of A−−H bond is about 335–355 kJ mol−1 (see, for instance [29, 30, 31]). Antioxidants belonging to this second family are called ‘chain breaking antioxidants’ or, more simply, ‘terminators’, for the sake of shortness. In the case of polyolefins, one can obtain an efficient stabilization with antioxidant weight fractions lower than 0.5%. In the case of polyene elastomers, antioxidants are, in general, used in concentrations higher than 1%. The most commonly-used method to evaluate the antioxidant efficiency consists in determining the oxidation induction time (OIT) at high temperature in pure oxygen (Figure 21.18). The efficiency of antioxidants depends also on several physical factors: 1 Their solubility into the polymer. If this latter is too low, the concentration threshold necessary to have an optimized stabilization cannot be reached. Numerous antioxidants are constituted of a long hydrocarbon chain, for instance thiodipropionates: S

CH2

CH2

C O

O

CnH 2n+1 2

with n = 12 (Dilauryl thiodipropionate (DLTDP)) or 18 (Distearyl thiodipropionate (DSTDP)).

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2 Their diffusivity into the polymer. If this latter is too high, their loss by migration can be too important and greatly reduce its efficiency. The solution consists in increasing their molar mass (hydrocarbon chain grafting, oligomer antioxidants, etc.) to reduce their diffusivity. 3 Their interactions with other additives. As an example, the efficiency of some antioxidants can decrease dramatically in the presence of fillers. Such antagonistic effects have been reported for associations of preventive antioxidants (hindered phenols or aromatic amines) and carbon black [33]. Some authors have attempted to provide an explanation. For some of them, the surface of fillers could catalyze the antioxidant oxidation [34]. For others, antioxidants could be neutralized by adsorption at the surface of fillers [35]. As another example, the efficiency of some antioxidants can increase strongly in the presence of another antioxidant. Such synergetic effects are well known for the past 50 years, for associations of preventive and terminator antioxidants, for instance for mixtures of sulphur + phenol, or phosphite + phenol. It is now currently used in industry (see, for instance [36]). For a long time, these synergetic effects have been interpreted in terms of chemical interactions between stabilizers or their by-products, for instance the regeneration of the phenol group by a reaction between the sulphur and phenol by-products [37], or more simply, by the fact that the sulphur reduces the hydroperoxide resulting from a stabilization step involving the phenol [39]. It can be easily demonstrated that, in fact, this synergetic effect results essentially from a simple kinetic effect due to the fact that each antioxidant acts in a different level of the radical chain process [28].

21.3.2.3

Hydrolytic Ageing

General Aspects Hydrolytic ageing results from chemical interaction between polymer and water. The main chemical event is a hydrolysis process which can be schematized as follows: ∼ A−−B ∼ + H2 O

→ ∼ A−−OH + H−−B ∼

In the most important practical cases, the broken chemical bond belongs to the macromolecular skeleton. Thus, each hydrolysis event is a chain scission. As a result, hydrolysis leads to polymer embrittlement at very low ratios of the reaction, like other chain scission processes studied previously, that justify the large amount of literature works dealing with chemical ageing of this type. A water molecule participates to the hydrolysis of a polymer chemical group if it penetrates into the sample. In this case, it modifies also the polymer properties and leads to a physical ageing. There are cases of ‘pure’ physical ageing by water absorption in the case of hydrophilic (but no hydrolyzable) polymers. On the contrary, there is no ‘pure’ chemical ageing, because each chemical ageing is followed by plasticization or swelling phenomena that are relevant to physical ageing. This ageing type is reversible (as long as the material does not damage), a drying restores initial material properties. Chemical ageing, on the other hand, is irreversible (Figure 21.19). Irreversible Hydrolysis Polycarbonate (PC), saturated linear polyesters (PET, PBT) and unsaturated polyesters (UP) crosslinked by styrene hydrolyze up to high conversion ratios, without observation of a significant slow down of reaction kinetics. For these polymers, one can consider, in a first approximation, that hydrolysis is irreversible. Thus: E + W → Ac + Al

(k)

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Figure 21.19 Shape of weight changes in humid (t < tS ) and dried atmosphere (t > tS ). (a) Case of ‘pure’ physical ageing: the weight gain tends towards an equilibrium value and, after drying, the material recovers its initial weight. Absorption and desorption rates are equal. (b) Case of hydrolysis (without extraction of small molecules from sample). Each hydrolysis event leads to a weight gain of 18 g.mol−1 . An irreversible sample weight variation is induced after drying. Note that the sample weight can increase, decrease or vary nonmonotonically according to chemical ageing type.

where E is an ester group, W a water molecule, and Ac and Al are respectively acid and alcohol groups resulting from hydrolysis. Each hydrolysis event leads to a chain scission (S). It can be thus written, in a nondiffusion controlled regime: dS = k[E] [W ] = k ([E]0 − S) [W ] dt

(21.40)

Since embrittlement occurs at a very low conversion ratio: S  [E]0 , hydrolysis does not modify much the polymer hydrophilicity: dS = k[E]0 [W ] ≈ cst dt

(21.41)

withK = k[E]0 [W ]

(21.42)

Hence: S = Kt

Let’s recall that the endlife criterion, corresponding to the ductile/brittle transition, is: MW = MF , and the corresponding critical number of chain scissions is: SF =

2 2 − MF MW 0

(21.12)

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Thus, the lifetime tF is: tF =

SF = K



2 2 − MF MW 0



1 k[E]0 [W ]

(21.43)

k obeys an Arrhenius law. [W] is determined from the water chemical potential. In the case of a slightly hydrophilic material: 1 if the material is immerged in pure water: [W ] = [W ] S where [W]S is the equilibrium value of water concentration into the material. 2 if the material is exposed to a humid atmosphere: [W ] = [W ] S

RH 100

where RH is the relative hygrometry. Thus, in the general case: tF =



2 2 − MF MW 0



100 1 1  [E]0 k0 exp −E RT [W ] S RH

(21.44)

with RH = 100 in the case of an immersion in pure water. The above equation provides a nonempirical relationship between lifetime and three factors representing the effects of polymer structure, temperature and environment composition. Reversible Hydrolysis In some cases, for instance PA11, hydrolysis is strongly autoretarded from the early period of exposure, and the sample weight tends towards an equilibrium value, that clearly indicates the existence of a reversible process: A+W

→ Ac + Am

(k H )

←

(k R )

where, in this case, A is the amide group, and Ac and Am are respectively acid and amine groups resulting from hydrolysis. One can write: dS = k H ([A]0 − S) [W ] − k R ([Ac]0 + s) ([Am]0 + s) (21.45) dt It can be easily shown that the equilibrium molar mass Mne (when t → ∞) is given by:  MW e = 2

kR k H [A]0 [W ]0

1/2 (21.46)

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Figure 21.20

829

Molar mass−temperature map (left). Lifetime versus temperature (right).

The molar mass is a decreasing exponential function of temperature (Figure 19.20). One can thus distinguish two distinct cases: 1 If T > TC , then MWe < MF . The material becomes brittle during its chemical ageing. 2 If T < TC , then MWe > MF . The material never becomes brittle, but reaches an equilibrium. Lifetime is theoretically infinite. In the case of PA11, TC ≈ 80 ◦ C.

21.4 Impact of Multiphase Structure on Ageing Processes Two aspects of multiphase structures are considered in this section: on the one hand, the semicrystalline nature of polymers implies the existence of amorphous regions and regions of well-packed molecular chains, generally of higher density, with interfacial regions between them. The second important family is that of composites, in which charges of variable nature and shape are incorporated. Generally, such charges are inorganic materials and are not specifically directly involved in ageing. We consider here how the multiphase nature of materials may affect ageing processes described above. 21.4.1

Structural Reorganization

Structural recovery concerns the amorphous phase of polymeric materials. In semicrystalline polymers, molecular motions within the amorphous phase are hindered to some extent by the presence of the crystalline phase. As a consequence, the glass transition temperature usually increases with the crystalline degree, and its manifestation (by DSC, for example) is smeared out. Therefore, structural recovery is a second order problem in this kind of material. The main changes in physical properties of semicrystalline polymers are therefore linked to variations of the crystalline degree itself. Such changes generally do not occur if the material is kept at a temperature below Tg . For higher temperatures, cold crystallization may occur, to an extent depending on the processing conditions of the material, in particular the cooling rate: a fast cooling rate freezes the material in a outof-equilibrium, disordered state. The molecular mobility prevents packing of macromolecules into crystal

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0.6 J/g/K

Tf = 258°C

Exo

Tg = 75.1°C

Tcr = 135.6°C 50

100

150 200 Temperature (°C)

250

Figure 21.21 Example of cold crystallization process observed by DSC on PET. The material was initially quenched to room temperature providing so essentially an amorphous phase. On subsequent heating, the glass transition and associated recovery of structural relaxation (small endothermic peak) are observed, followed by cold crystallization at 135 ◦ C and melting at 258 ◦ C.

structure. With time and/or in combination with a temperature increase, so-called cold crystallization occurs. This marginally concerns materials like polyethylene which crystallize very fast and can virtually not be obtained in an amorphous state. Materials like Poly(ethylene terephthalate) or Poly(ethylene naphthalate) can easily be obtained either in the amorphous or crystalline structures and are therefore prone to cold crystallization (cf. Figure 21.21). Negative impact (that can be viewed as positive effects depending on the targeted application) of a crystalline degree increase can be as follows:

r r r r r

increase in brittleness of the material loss of elasticity, increase in Young’s modulus volume reduction due a crystalline phase more dense than the amorphous one, leading e.g. to void formation in the bulk, shrinkage, and decohesion from charges in a composite or from adjacent pieces change in material aspect (from transparent to translucent due to diffusion of light by crystallites) apparition of mechanical and dielectric losses due to the so-called α-relaxation process (see e.g. [39] for the outlines of such a process).

A peculiarity of shrinkage effects in oriented polymeric films is worth mentioning here. Polyvinylidene fluoride (PVDF) is known for its ferroelectric properties. Forty years after the discovery of piezoelectricity in PVDF, it remains one of the best piezoelectric polymers available so far. For inducing ferroelectric properties to PVDF, uniaxial or biaxial mechanical drawing of films, to a ratio of the order of 4, has to be carried out, inducing change in the crystalline phase from a nonpolar α-phase to a polar β-phase. A high electric field is then applied to provide a macroscopic polarization to the films. The point is that when the material is heated to temperature of ca. 70 ◦ C and above, shrinkage of the films occurs, implying motion of the crystalline regions. This results in a lowering of the macroscopic polarization and henceforth an ageing of the pyroelectric and piezoelectric properties [40]. The copolymerization of VDF units with TrFE (trifluoroethylene) produces directly crystallization in the polar phase, so avoiding drawing of the films and improving the thermal stability of the electroactive properties of the materials.

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Figure 21.22 Example of a transcrystalline region formed around a fiber during crystallization of a PPS/carbon fiber composite.

In composite materials, structural recovery is concerned, to a first approximation, in the same way as for amorphous materials if micrometric scale particles are considered. In nanocomposites, i.e. composites containing nanometer scale particles, it can be anticipated that hindering of molecular motions occurs and impedes such effects, as it does for semicrystalline materials. Crystallization mechanisms can be significantly modified by the presence of charges. Indeed, two processes of crystallization germination are at play, which are homogeneous germination (self ordering of macromolecules acting as a core of the crystal or spherulite) and heterogeneous germination, i.e. crystal growth starting from a defect or from small crystals in surfusion. Therefore, charges themselves may act as ‘defects’ inducing crystal growth preferentially from the surface of such defects. Figure 21.22 shows an example of heterogeneous germination of Poly(phenylene sulphide) –PPS- on a Carbon fiber: very specific morphology, called transcrystalline region, is obtained in this case [41]. It results from a dense crystal nucleation on the fiber and a growth of crystalline lamella along radial directions of the fiber. It has been reported that the interface between the transcrystalline region and material bulk can constitute a weak point from the mechanical point of view in fiber-reinforced semicrystalline polymers.

21.4.2

Diffusion Controlled Processes

Generally speaking, diffusion processes tend to be slowed down when particles are incorporated into a matrix. Already geometrical considerations can support the statement. As diffusion proceeds through free volume occupancy, the presence of crystalline lamellae and/or of inorganic particles reduces the volume effectively available for diffusion since it concerns essentially amorphous phases. In addition, molecular motions are hindered by the presence of charges/crystallites and hence it can be anticipated that the diffusion coefficient for small molecules be lower in the amorphous phase of a semicrystalline material than in a purely amorphous material. There are, however, notable exceptions to this statement, particularly as regards water. In epoxyglass composites as an example, water tends to accumulate at the interfaces between matrix and filler and to form so-called water shells. If at this stage, chemical degradation is not discussed, the presence of water at such interfaces can have tremendous consequences, like a huge increase in dielectric losses, along with the formation of weak cohesive points. This is a critical issue for the mechanical resistance of the composite as it is expected that stresses are transferred between the two phases through this interface. Except for this

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specific case of water in epoxy-based composites which is dealt with specifically in the next section, it can be considered that composites are detrimental to diffusion, which can be used as a positive property.

21.5 Practical Impact of Physical Ageing on Use Properties So far, we have dealt with ageing mechanisms in polymer and composites that appear under ‘normal’ ambient conditions, essentially as a combination of humidity, oxygen and temperature. In the following, we review issues associated with material evolution under functional stresses, i.e. specific environmental constraints that materials are supposed to sustain during the life of systems with which they are incorporated. In many instances, the impact of ageing mechanisms on mechanical properties has been discussed in the previous sections. Therefore, we shall consider only briefly these aspects therein. Given the extremely broad range of applications in which polymers are incorporated today, it is impossible to build an exhaustive review of those stresses. Instead, we propose to consider electrical insulating properties, an important application to our eyes in which the ‘multiphase’ nature of materials is exploited. The issues associated with the ‘multiphase’ nature of materials are discussed in relation to the demanding performances for such application. 21.5.1 21.5.1.1

Water-Induced Mechanical Damages in Composites Osmotic Cracking

In the case of composites made of glass fibers and unsaturated polyester matrix (boat hulls, swimming pools, etc.), weight changes due to water absorption can present the peculiar shape illustrated in Figure 21.23. This behavior can be explained as follows: 1 The initial stage corresponds to water absorption by the classical dissolution−diffusion mechanism. At time t1 , the material reaches its sorption equilibrium. 2 However, water reacts chemically with the polymer. Since this latter presents a lot of chain ends, chain scissions close to chain extremities lead to an accumulation of small organic molecules (monomers, oligomers). When these small molecules (which are more polar than the polymer) reach their solubility threshold into the polymer, they induce, at time t2 , a phase separation and generate micropockets of highly hydrophilic liquid. The water entrance in these micropockets generates an osmotic pressure and,

Figure 21.23

Shape of weight changes for a material composite subjected to an osmotic cracking.

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thus, a stress state which can initiate cracks at the micropocket surface. Between t2 and t3 , the sample weight increases because cracks propagate and receive more and more water. 3 At time t3 , cracks coalesce and the solutes leave the sample, producing a quick drop in sample weight. In the case of composite materials, cracks develop in sub-layers of the material and propagate parallel to the sample surface, which leads to blister formation. Material ability to generate blisters in the presence of water is essentially linked to the concentration of small molecules accumulating in the matrix. This concentration depends on three factors: 1 the concentration of initially present small molecules (in particular, the initiator of polymerisation and its reaction products) 2 the initial concentration of chain ends (all the more so high since the molar mass of initial linear unsaturated polyester is low) 3 the rate of polymer hydrolysis (leading to new chain ends). The effect of these different factors has been checked experimentally (for instance, [42]). 21.5.1.2

Interfacial Hydrolysis

In some cases of composite materials based on matrices particularly resistant to hydrolysis (vinylester, nonhydrolyzable thermoplastics), humid ageing leads to an alteration of mechanical properties which cannot be attributed to a chemical interaction between water and matrix. Microscopic or mechanical analyses (for instance, resistance to interlaminar shearing) show the existence of a fiber–matrix debonding. Moreover, the composite stability, in given conditions, depends closely on the nature of the coupling agent. Thus, hydrolysis reactions take place at the fiber–matrix interphase. There are numerous experimental methods to support qualitatively this phenomenon. But, to our knowledge, there are no methods allowing quantifying, in an undisputable way, the numbers of chemical events responsible for this evolution. At the present time, there is no possibility to perform a fine kinetic analysis of this chemical ageing type. 21.5.2

Ageing of Electrical Insulations

Polymers are involved as electrical insulating materials in all the systems of electrical engineering, electrotechnics and electronics. Some examples of application and of materials used as insulations are given in Table 21.5. Criteria behind the selection of a given material depend on the intrinsic electrical performances expected (see below), on the resistance to other environmental stresses (mechanical, UV, temperature, humidity. . .), on the processability of the material and on economical factors. In the following, we detail the expected electrical properties and consider their evolution in bulk polymers and composites. 21.5.2.1

Critical Electrical Properties

Field Strength The most important electrical property for an electrical insulation is the field strength, i.e. the maximum electrical field the material can sustain without breakdown. Typical values are in the range 15 to 600kV/mm, depending on materials, on the stress form (DC, AC, impulse voltage) and on sample thickness. The property must be maintained during the service life of the material. The monitoring of field strength is often a way of evaluating, from the electrical standpoint, the resistance of material to degradation during

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Table 21.5 Field of application of some polymers and composites in Electrical Engineering. Electrical Engineering – Electrotechnics

Electronics – Power electronics

Application

Materials

Application

Material

High voltage cables

-Polyethylenes, PVC -Carbon black reinforced polymers

Capacitors

Outdoor insulations

Epoxy, silicone

Power transformers

Silica-filled epoxies

Low voltage cables

XLPE, EPR, PVC, ETFE (fluorinated copolymers), silicones Epoxydes, polyimides, polyurethanes, polyamides Epoxy/mica composites

Components packaging Power electronics insulation Circuit pressboards

Polypropylene, PET, polycarbonate, polystyrene, fluorinated polymers, polysulfones Silicone gels

Wires: varnishes

Bus bars

Polyimide Fiberglass reinforced epoxies

accelerated ageing tests. Usually, the field strength is approached as a short-term process, i.e. ageing effects associated with the application of the stress are not considered. Three mechanisms of dielectric breakdown are worth mentioning here. These are: avalanche breakdown, electromechanical breakdown and thermal breakdown. Avalanche multiplication occurs at high field strengths when carriers acquire sufficient kinetic energy between collisions (scattering events) with the matrix to give a high probability of ionization, with the generation of more carrier. Thus a small current can be greatly multiplied by an increase in the number of carriers until it is sufficiently large to cause irreversible damage. This type of mechanism has been used to explain breakdown in oxidized films. Thermal breakdown occurs when the heat input cannot be balanced by the heat losses from the insulation either macroscopically, or more usually, in a small area. As power is dissipated by the insulation, heating occurs which usually causes an exponential increase in the electrical conductivity as more carriers become available for conduction. If the electrical stress is maintained, the current density increases in the area at elevated temperature. This serves to further increase the local temperature through Joule effect and hence the conductivity, and so thermal runaway may occur. Electromechanical breakdown occurs due to the electrostatic attraction of the electrodes which decreases the width of the insulation by an amount depending on the Young’s modulus. If the applied voltage is maintained, the field increases due to the decrease in thickness, thereby increasing the attraction force (which varies as a square function of field) further. Aside from the mechanism itself, the form of breakdown, either homogeneous or filamentary, is questioned. Contrarily to mechanical breakdown which requires extensive damage to definitively occur, electrical breakdown can be initiated and accomplished through only a weak channel through the insulation. In this respect, models have been developed for filamentary breakdown [4] of either a thermal or electromechanical nature. For the latest mechanism, the filamentary process appears much more likely than homogeneous ones. The mechanism has analogy with fracture mechanics in which a crack spontaneously propagates according to the Griffith’s criterion. It is substantiated in some works by the correlation between field strength and

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Young’s modulus for e.g. experiments at variable temperature [44]. Filamentary processes could be initiated from weak points in the bulk or at the surface of the material. Breakdown statistics tend to show that the probability of having a breakdown event increases with the volume of small. This tends to provide support to defect-induced initiation of breakdown as opposed to homogeneous processes. The above processes are general and are not specific to polymers. Dielectric breakdown processes in polymers are complicated due to their highly complex structures. Therefore, much work has been devoted to clarifying the correlation between the breakdown characteristics and the properties inherent to polymers. Since condensed states of polymers change with temperature, the temperature dependence of their dielectric breakdown is of prime importance in analyzing their breakdown mechanisms. Roughly, it can be considered that breakdown at low temperature in polymers could be related to avalanche breakdown, basically because the two other processes become less probable. Thermal and electrical breakdown would be favored for high conductivity materials and soft materials, respectively. Electrical Ageing We refer here to ageing processes in which the electrical stress is a necessary condition to ageing. All the cited mechanisms are expected to lead to premature breakdown in the long term. The more severe forms of degradation in this respect are water trees, partial discharges and electrical trees. Water trees are of concern where the main application is submarine cables. It results from the application of an electrical stress combined with contact with an aqueous electrolyte. Water trees do not generally form interconnected channels. After the initiation stage, trees grow very rapidly at first and then the speed of growth progressively decreases. These features have been tentatively explained by a great number of mechanisms of electromechanical, electrochemical oxidation or electro-mechano-chemical nature, but in spite of this diversity none of the mechanisms are adequate to describe all the basic observations. The common scheme is that of a microvoid formation (initiation) followed by extension of that void along a privileged direction. Electrical trees are generated in field-enhanced regions of the material that can be a void or a metallic conducting center (the typical experiment for observing it being a needle electrode). Electrical trees consist of connected channels (hollow tubules) a few μm in diameter, with branches tens of microns long. The walls of the channels can be carbonized and are only weakly conducting. They are not considered as intrinsic runaway breakdown mechanisms, rather a cumulative damage of the insulation leading to an enhancement of the failure probability. Finally partial discharges are discharges occurring in the gas included in voids of some tens of μm size included in the polymer. Erosion of the internal surface of the voids by particles from the discharges leads to an enlargement of the void and eventually formation of trees. In practice, the actual life of the insulation is actually not controlled by the extension of such voids, but by the time for partial discharges to be detected. Figure 21.24 provides a sketch for a possible ageing and breakdown scenario for a material free from preexisting voids. Considerable improvement in the reliability of high voltage cables has been obtained by improving the quality of products and, notably, in avoiding the presence of microvoids inside insulations. Actually, the fundamental question for electrical ageing of insulations is the nature of mechanisms leading to the formation of voids of size large enough to sustain a discharge process [45]. This criterion remains one of the most important quality control parameters in the production of such cables. Acceleration of electrons in regions of lower density in the material (free volume regions) followed by impact excitation on molecules is one way through which molecular-excited states can be produced and may lead to chemical degradation [46]. Along these lines, the actual reactivity of centers’ localizing electrical charges (space charge) and the electrostatic and electromechanical energy stored around such centers are other lines of thought for explaining such phenomena. Space Charges In the absence of voids, space charge-induced phenomena represent one of the most important lines of research regarding electrical ageing of insulation [47]. Space charges are electronic or

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Charge or field-induced ageing

Void-free material Degradation reactions – molecular scale SC-induced phenomena Excited states due to hot e- or charge recombination

Nano-void Growing of voids SC-induced mechanical forces Increasing mean free path of carriers in LD regions Solid-gaz interface: hot e - in the gas

Discharges-induced degradation

Micro-void Erosion of cavity boundary by PD Chemical degradation due to the discharge Hot e- , UV, reactive species, T effects…

Electrical tree Propagation Hot e- from plasma age the dielectric at tree tip Electromechanical forces

Breakdown Figure 21.24 Sketch of a possible breakdown scenario in a void-free insulation material. Hot e- stands for hot electrons, i.e. electrons of kinetic energy of a few eV, able to excite molecules by impact. PD stands for partial discharges.

ionic carriers present in an insulation. Depending on their trapping energy, they can move under the effect of the electric field. Numerous techniques of detection of such charges have been developed during the last 30 years. The negative impact of space charges can be viewed from different standpoints:

r r r

the acceleration of such charges may give rise to excitation of molecular centers as explained above recombination of charges of different polarity is another way of producing excited states in the insulation the accumulation of charges in some regions implies, in virtue of the Poisson’s equation, changes in the actual electric field in respect of the design field. It implies necessarily regions of enhanced electric field that may induce premature breakdown of the insulation.

The last point is a critical issue for high voltage DC applications where, contrary to AC cases, there is no extensive experience regarding synthetic insulations, in particular for cables. Surface Discharges Surface properties, for example, are essential features for outdoor insulations. Pollution of the surface leads to an increase of the conductivity, surface discharges processes (tracking discharges) and flashover. In a number of applications the surface has to be resistant to electrical discharges, in particular avoiding erosion of the material. The environmental stresses are generally involved in failure associated with tracking discharges. That factor is often moisture that condenses on degraded or contaminated surfaces making them conduct current. Discharges occur when the current path is broken by drying. The discharge may cause local carbonization.

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More wetting causes further discharge, more carbonization, and finally, failure. The intensity of this effect of moisture is so dependent upon the surface condition that it is often separated from the aging mechanism and considered simply as a mode of failure [48]. Besides, moisture can contribute to a dominant mode of degradation that is very different from hydrolytic attack. When the environment includes rain, fog or any other form of condensate moisture, as well as voltage, a special type of deterioration occurs. It is the same process as in the failure mechanism, but now it determines the life and not simply the end of life, and the details are more important to understand. Leakage current flows across the surface and the heat generated causes a drying of the moisture film. A circumferential ‘dry-band’ forms, which partly interrupts the charge flow. Most of the applied voltage is supported by this dry-band, and arcs jump across it. These arcs are a type of discharge variously called ‘arclets’ or ‘scintillation’. They are fed by the current through the moisture film, so that they are hotter than partial discharge would be. They impinge on the insulation surface, causing an ambient of fast moving, transient high temperature spots, each lasting a few ms. While they last, the atmosphere in the area of the arcs is an ionized plasma with a temperature of thousands of degrees. The surface of the insulation will react chemically with the constituents of the plasma, and degrade with the heat, causing a spot of carbon or other conductive residue to form on the material surface. This localized degradation may become an attachment point for the arc terminus (the cathode spot) so that the rate of degradation becomes self-accelerating. The degradation zone extends itself from one electrode to the other, and the former insulator ceases to be an insulator. This is one possible way to view ‘surface tracking’. Essential features that must be maintained during the life of outdoor insulation are surface hydrophobicity in respect to water effects, and resistance to sunlight as a general requirement. Other Issues A number of other electrical properties can be important for insulations and dielectrics, depending on the application. Among them, bulk dielectric properties, encompassing conductivity, permittivity and dielectric losses are fundamental parameters for such materials. Changes that can be observed here are induced more by environmental parameters such as humidity and water uptake, UV exposure, thermal stresses than by pure electrical ageing. In the following, we give an overview of the impact of multiphase systems, being semicrystalline polymers and composites, on the properties described above.

21.5.2.2

Ageing of Electrical Properties in Multiphase Systems

Polymers in General Any ageing process able to produce voids within the insulation does have negative impact on the electrical performance. On the one hand, voids may initiate breakdown processes as explained above. On the other hand, diffusing species, among which water may segregate in such voids and produce therefore a composite material. Dielectric losses can be induced by such processes due to the difference in permittivity of the two phases. Such void-containing regions can be induced by mechanical stresses along with other chemical ageing processes producing crack. An incomplete cure of a resin can also form regions of low density. Post-curing of a resin produces volume contraction and is also able to produce voided materials; the same may happen with thermal cycling due to differences in thermal expansion coefficients. Voltage stabilizers, i.e. additives for improving water and electrical treeing resistances and the electrical strength in polyethylene cables have been known for several decades but have not been widely used. The principles behind the stabilization mechanism are outlined as follows [49]. Stabilizers are generally aromatic groups being more easily ionized than the polymer chain, forming a stabile radical cation when an excited electron is captured. This means that rather than dissipating its energy by breaking bonds in the polymer chain the excited electron reacts with the stabilizer. This cation radical of the stabilizer will in turn be able to abstract

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energy from another excited electron or a radical center on the polymer, leaving them unexcited. When these steps have taken place the voltage stabilizer has been regenerated to its original state and no depletion of the stabilizer due to consumption will occur. The stabilization effect has been correlated to the ionization potential of such groups. However, the low solubility of the aromatic voltage stabilizers in polyolefins has been a dilemma. One of the main problems with voltage stabilizers until quite recently has been the depletion of stabilizer due to migration from the material, and hence a decrease in time of its efficiency. Alternative routes, such as grafting stabilizers on the chains [50] or design of molecules with higher stabilization efficiency have been proposed [51]. Thermal degradation generally goes with an increase in the conductivity of the material and is therefore detrimental to the insulation. The same also holds for UV ageing, particularly for surface insulation. Cracking, hardening and stresses cycling may induce decohesion between metallic parts and the insulation. From there, partial discharge can induce triggering and so erosion of the material. Semicrystalline Polymers In some reports on the impact of microstructure on breakdown phenomena, it has been shown that spherulites may represent weak points of the insulation: discharge channels appearing along spherulites boundary [52]. This not completely surprising as the interfaces between spherulites are regions of lower density, that may also contain a high concentration of defects. It can be suspected that the current density is higher in these regions, so producing damage. Crystalline regions constitute in general regions of lower mobility for charges. This, along with the fact that molecular motion is hindered so that free volume is reduced, leads to a lower conductivity for a semicrystalline polymer. The same holds for permittivity: the latter exhibits a significant step across the dielectric relaxation associated with the glass transition. An increase in crystallinity lowers the dielectric permittivity which can be detrimental for application as capacitors (energy storage). The counterpart is a decrease in the losses associated with the β-relaxation. An important consequence of the semicrystalline nature of polymers is the appearance of dielectric relaxation processes associated with paracrystalline regions (α−relaxation), at temperatures above Tg . Composites There are numerous examples where, aside from economic benefits, composites bring a positive impact on insulation performance as a whole. Essentially, the gain obtained concerns thermal and mechanical properties, including a decrease in the thermal expansion coefficient and an improvement of the thermal conductivity and resistance to surface discharges. The latter effect can be understood by the strong resistance of inorganic particles (e.g. mica in transformer bars) to discharges, forming a protective layer and avoiding erosion. In outdoors insulations, silicone elastomers are used more and more with silica or ATH (alumina trihydrate) particles as fillers, one of the advantages of silicones being the possibility for low molecular weight polymeric chains to diffuse up to the surface of the insulation, so ensuring a regeneration of the surface properties (hydrophobicity recovery, incorporation of pollutants into the insulation bulk). See [53] for a review of ageing and recovery of silicone-based composites in the field of outdoor insulations. There are, however, negative effects of the incorporation of fillers into composites. As with silicones, coatings develop leakage currents all the earlier that the filler content is high. Also, the incorporation of particles tends to decrease the hydrophobic properties. In silica-filled epoxy materials, the dielectric strength is divided by a factor of two in respect to the pure resin for 50% load. One of the reasons is that the particles are incompletely dispersed in the resin and therefore provide weak points. Also, water uptake can be important (of the order of 2%), implying a tremendous increase in dielectric losses. The effect is much more pronounced in composites than in pure epoxy and is further increased in nanocomposites due to the large specific surface of particles [54, 55]. Simple calculations indicate that the water is not dissolved in the resin but aggregates around silica particles forming all sorts of interfacial layers, and possibly decohesion.

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The last example concerns printed circuit boards (PCB), usually made of glass fiber-reinforced epoxies. Important issues are currently coming in relation to the increase in working temperature and voltage induced notably by the development of power electronics and by the trend towards high density integration of electronic circuits implying higher electric field applied to the insulation layer of a PCB. Aside from the surface effects that have been addressed for a long time, like oxidation of metallic layers and the formation of whiskers, new features have been revealed such as ionic migration in the bulk of the PCB [56] and possibly also migration of metal atoms and electrochemical processes along glass fibers from nearby contacts. Such effects may lead to short-circuit and failure of the device. The above examples of ageing in polymer and composites used for their electrical insulating properties are in no way exhaustive and have been introduced for illustration purposes. Given the wide spread of application fields, of environmental conditions, of working stresses, of geometry and of nature of materials, weak points in terms of properties can be extremely variable. Concerning the relationships between the intrinsic electrical properties and the composite nature of materials, generally, the introduction of fillers brings a much more complex behavior (e.g. dielectric losses, interfacial effects), weaknesses (e.g. electrical strength, extra ageing routes, sensibility to water) and some improvements (erosion resistance under discharges, UV-protection). Nanocomposites have certainly positive effects in respect to microcomposites in the sense that they tend to maintain the advantages brought by filling the material and to attenuate drawbacks. Even though intrinsic properties are valuable, their development in this field is still dependent on reliability assessment, on processing constrains, availability and economical considerations.

21.6 Concluding Remarks The purpose of this chapter was to present the main polymer ageing types likely to take place in practice and, in particular, to provide the reader with important pieces of information such as: 1 There is not a single ageing type, but several ageing types susceptible to superimpose and interact. 2 Ageing obeys a more or less complex set of physical and chemical mechanisms, each one obeying a well-known kinetic law. 3 In the past half century, researchers have made many efforts to elucidate these elementary mechanisms. Now, the main challenge is the elaboration of nonempirical kinetic models simulating satisfyingly the kinetics of such complex ageing processes (for instance, chemical ageings involving several dozens of chemical reactions, kinetically controlled by reactant diffusion, etc.) and their consequences on polymer structure, and physical and mechanical properties at the pertinent scales: the molecular, the macromolecular and the macroscopic scales; 4 Problems of thermal and radio-oxidation, and hydrolysis, kinetically controlled or not by diffusion reactant, are now solved nonempirically using numerical models derived from realistic mechanistic schemes (respectively, [57, 58]). However, significant efforts remain to be made in the field of polymer photo-oxidation and interphase hydrolysis in composites materials, stacked assemblies, etc. The problem of interactions between oxidation and humid ageing is also interesting. Indeed, on the one hand, some oxidized groups can be highly polar (for instance, alcohols and acids). They can therefore modify substantially the water transport properties into a thick polymer piece. On the other hand, other oxidized groups (for instance, esters and amides) can hydrolyze easily. 5 Actual ageing phenomena depend on the environmental stresses acting on materials. Besides the skills in materials chemistry needed to apprehend the stability range of materials, there are fields like electrical ageing that are still unsatisfactorily explored when dealing with long-term ageing, particularly for multiphase and composite systems.

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References 1. Colin X., Audouin L., Verdu J., Polym. Deg. Stab. 92(5), 886–897, 898–905 and 906–914 (2007). 2. Colin X., Audouin L., Verdu J., Rozental M., Rabaud B., Martin F., Bourgine F., Polym. Eng. Sci. 49(7), 1429–1437 and 49(8), 1642–1652 (2009). 3. Struik L.C.E., Physical Ageing in Amorphous Polymers and Other Materials, Elsevier, Amsterdam, 1978. 4. Tiemblo P., Guzman J., Riande E., Mijangos C., Reinecke, H., Polymer 42, 4817–4823 (2001); ibid 8321. 5. Bernier G.A., Kambour P., Macromolecules 1, 393 (1968). 6. Mc Kague E.L., Halkias E., Reynolds J.D., J. Compos. Mater. 9, 2 (1975). 7. Shen C.H., Springer G.S., J. Compos. Mater. 10, 2 (1976). 8. Whitney J.-M., Browning C.E., ASTM STP 658, 33 (1978). 9. M¨uller-Plathe F., Polymerica Acta 45, 259–293 (1994). 10. Masaro L., Zhu X.X., Prog. Polym. Sci. 24, 731–775 (1999). 11. George S., Thomas S., Prog. Polym. Sci. 26, 985–1017 (2001). 12. Wypych G., Handbook of Plasticizers, ChemTec Publishing, Toronto, Canada, 151–170, 2004. 13. Wypych G., PVC Degradation and Stabilization, ChemTec Publishing, Toronto, Canada, 2008. 14. Saito O., J. Phys. Soc. Jap. 13(2), 198 (1958) and 13(12), 1451 (1958). 15. Fayolle B., L. Audouin, J. Verdu, Polym. Deg. Stab. 70(3) 333 (2000). 16. Bellenger V., Verdu J., J. Appl. Polym. Sci. 28, 2599–2609 and 2677–2688 (1983). 17. Bauer D.R., Mielewski D.F., Gerlock J.L., Polym. Deg. Stab. 38, 57–68 (1992). 18. Griffith A.A., Phil. Trans. Roy. Soc., Series A 221, 163–198 (1920). 19. Audouin L., Verdu J., in Radiation Effects on Polymers, R.L. Clough and S.W. Shalaby (Eds), ACS Symposium Series 475, American Chemical Society, Washington D.C., p. 473, 1991. 20. Bolland J.L., Gee G., Trans. Faraday Soc. 42(3-4), 236 and 244 (1946). 21. Kamiya Y., Niki E., in Aspects of Degradation and Stabilisation of Polymers, H.H.G. Jellinek (Ed.), Elsevier, New York, p. 86 (1978). 22. Korcek S., Chenier J.H.B., Howard J.A., Ingold K.U., Can. J. Chem. 50, 2285 (1972). 23. Coquillat M., Verdu J., Colin X., Audouin L., Nevi`ere R., Polym. Deg. Stab. 92(7), 1334–1342 (2007). 24. Gillen K.T., Wise J., Clough R.L., Polym. Deg. Stab. 47, 149 (1995). 25. Colin X., Verdu J., Plastics Rubber Comp.: Macromol. Eng. 32(8/9), 349–356 (2003). 26. Assadi R., Colin X., Verdu, J., Polymer 45(13), 4403–4412 (2004). 27. Colin X., Richaud E., Verdu J., Monchy-Leroy C., Rad. Phys. Chem., 79(3), 365–370 (2010). 28. Verdu J., Rychly J., Audouin L., Polym. Deg. Stab. 79(3), 503 (2003). 29. Bordwell F.G., Zhang X.M.A.N., J. Phys. Org. Chem. 8(8), 529 (1995). 30. Mulder P., Saastad O.W., Griller D., J. Amer. Chem. Soc. 110(12), 4090 (1988). 31. Zhu Q., Zhang X.-M., Fry A.J., Polym. Deg. Stab. 57(1), 43 (1997). 32. Schwarzenbach K., Gilg B., M¨uller D., Knobloch G., in Plastic Additives Handbook, 5th edition, H. Zweifel (Ed.), Hanser, Munich, Switzerland, 2000. 33. Shelton R.J., in Polymer Stabilization, W.L. Hawkins (Ed.), Wiley-Interscience, New -York, p. 104, 1971. 34. Hawkins W.L., Worthington M.A., J. Polym. Sci. 62(174), S106 (1962). 35. Kuzminskii A.S., Lyubchanskaya L., Khitrova N., Bass S., Rubber Chem. Technol. 26, 858 (1953). 36. Zweifel H., Plastics Additives Handbook, 5th edition, Hanser, Munich, Switzerland, 2000. 37. Hawkins W.L., Worthington M.A., J. Polym. Sci. Part A 1(11), 3489 (1963). 38. Pospisil J., Nespurek S., in Handbook of Polymer Degradation, S. Halim Hamid (Ed.), Marcel Dekker, New York, p. 259, 2000. 39. Teyss`edre G., Grimau M., Bern`es A., Martinez J.J., Lacabanne C., Polymer 35, 4397–4403 (1994). 40. Teyss`edre G., Lacabanne C., Ferroelectrics 171, 125–144 (1995) 41. Boudet A., Teyssedre G., Caramaro L., Chabert B., Bull. Soc. Chim. Belg. 98, 747–750 (1989). 42. Gautier, L; Mortaigne, B; Bellenger, V., Composites Sci. Technol. 59(16), 2329–2337, 1999. 43. Fothergill J. C., IEEE Trans. Electr. Insul. 26, 1124–1129 (1991). 44. Kim H.K., Shi F.G., IEEE Trans. Dielectr. Electr. Insul. 8, 248–252 (2001).

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Ageing and Degradation of Multiphase Polymer Systems 45. 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57.

58.

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Morshuis, P.H.F., IEEE Trans. Dielectr. Electr. Insul. 12, 905–913 (2005). Laurent C., Teyssedre G., Nucl. Instr. and Meth. in Phys. Res. B 208, 442–447 (2003). Teyssedre G., Laurent C., IEEE Trans. Dielectr. Electr. Insul. 12, 857–875 (2005). Starr W.T., IEEE Trans. Electr. Insul. 25, 125–136 (1990). Englund V., Hjertberg T., Huuva R., Gubanski S.M., Proc. 2007 International Conference on Solid Dielectrics, Winchester, UK, July 8–13, 203–206 (2007). Peruzzotti F., Martinotto L., Del Brenna M., in Cable, in particular for transport or distribution of electrical energy and insulating composition, US Patent 6696154, 2004. Englund V., Huuva R., Gubanski S.M., Hjertberg T., Polymer Degr. Stab. 94, 823–833 (2009). Sawa G., IEEE Trans. Electr. Insul. 21, 841–846 (1986). Reynders J.P., Jandrell I.R., Reynders S.M., IEEE Trans. Dielectr. Electr. Insul. 6, 620–631 (1999). Zou C., Fothergill J.C., Rowe S. W., IEEE Trans. Dielectr. Electr. Insul. 15, 106–117 (2008). Adohi I.P., Guillermin C., Rain P., Rowe S.W., 2004 Annual Report Conference on Electrical Insulation and Dielectric Phenomena, 158–162 (2004). Natsui M., Echigo Y., Tanaka T., Ohki Y., Maeno T., Proc. 2008 International Symposium on Electrical Insulating Materials, Sept. 7–11, 2008, Yokkaichi, Mie, Japan, pp. 376–379 (2008). Colin X., Fayolle B., Audouin L., Verdu J., in Advanced in Chemistry Series 1004: Polymer Degradation and Performance, M.C. Celina, J.S. Wiggins and N.C. Billingham (Eds), Oxford University Press, Washington, pp 121–134, 2008. Jacques B., Werth M., Merdas I., Thominette F., Verdu J., Polymer 43, 6439 (2002).

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22 Fire Retardancy of Multiphase Polymer Systems Michel Ferriol LMOPS, Universit´e Paul Verlaine Metz, Saint-Avold, France

Fouad Laoutid Materia Nova asbl, Mons, Belgium

Jos´e-Marie Lopez Cuesta CMGD, Ecole des Mines d’Ales, Ales, France

22.1 Introduction Polymeric materials are used more and more in our environment due to their main properties: low weight, easy processing, etc. However, they are also easily flammable, their combustion leading to the production of corrosive or toxic gases and smoke. The consequences of fires are dramatic. Every day in Europe 12 people are victims of fires and 120 are severely injured. About 80% of all these deaths occur in residential buildings. The total economic damage can be estimated at about € 25 billion per year. So, improving the fire retardant behavior of polymers is a major task for improving fire safety. Numerous compounds and formulations for improving polymer fire resistance can be found in the literature. Among them, brominated compounds show a high efficiency and three chemical groups of brominated compounds dominate the current scene: polybromo diphenyl ethers (PBDE), hexabromocyclododecane (HBCD) and brominated bisphenols (particularly TBBP-A). However, they present strong drawbacks: release of toxic degradation products, bio-accumulation, and persistence, particularly in aquatic and marine environment. Hence, regulations in many developed countries have phased out some brominated FRs, such as the WEEE Directive banning PBDE in electric and electronic equipments in the European Union. In addition, with fire safety requirements becoming more and more rigorous, a new challenge arises for researchers to develop new efficient, halogen-free and environmentally-friendly flame retardant systems for polymer materials. Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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Among several others, a promising solution consists in the association of mineral fillers, and particularly nanofillers, with conventional fire retardant compounds. In this scheme, the aim of this chapter is to give to the reader some keys to the understanding of fire behavior in polymer materials and how it can be improved and controlled. The chapter is divided into four parts: 1. 2. 3. 4.

the fundamentals of combustion and flame retardancy of polymers the tests used to characterize fire behavior the most common flame retardants and how they work synergism in combinations of (nano)fillers and conventional flame retardants.

22.2 Combustion and Flame Retardancy of Polymers 22.2.1

Combustion of Polymers

A polymer is a material with a high concentration of carbon and hydrogen, chemical elements well known for their reducing properties. Therefore, it contains a significant fuel load, releasing large quantities of energy if the necessary conditions for ignition are met. The combustion takes place when the ‘fire triangle’ is on (Figure 22.1) which requires the combination of three elements: a fuel (the polymer), a combustive agent (the oxygen of air), and a heat source. Prior to the ignition, the polymer undergoes overheating due to the temperature increase in the vicinity of the heat source. During this first stage, thermoplastic materials soften or ‘melt’ whereas thermosets are not affected due to their crosslinked structure. At equal irradiation, the evolution of the material depends on its thermal properties: heat conductivity, thermal diffusivity, heat capacity, melting enthalpy, and so on. In a second stage, when the temperature reaches a crucial value, the polymer begins to degrade. The thermal degradation generates molecules of low molecular weight giving combustible gases or not. It usually occurs by three mechanisms that can possibly enter in competition: depolymerization, random scission of the polymer chain, degradation of side groups [1]. The degradation temperature is evidently linked to the bonding energy, the weakest chemical bonds being the first to break. Then, the combustion occurs as follows: ignition of combustible gases and flame propagation in a very exothermic radical mechanism in which HO◦ and H◦ radicals, so-called ‘hot radicals’, play a leadership role and lead to combustion products (CO, CO2 , soots, toxic gases or not, . . .). The heat generated at this level and

Heat

Fuel

Combustive

Polymer

Oxygen

Figure 22.1

Representation of the fire triangle.

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gaseous phase

non combustible gases

O2

COMBUSTION

combustible gases

845

combustion products

heat transfers

O2

thermal degradation condensed phase

Figure 22.2

polymer Schematic diagram of combustion of polymers.

the ensuing heat transfers (by convection, conduction and radiation) maintain the degradation of the polymer and thus feed the flame in gas fuel (Figure 22.2). The phenomenon is therefore self-sustaining.

22.2.2

Flame Retardancy

Reducing the risk of ignition in case of contact with a heat source or deceleration of combustion if the polymer or a neighboring material has already caught fire requires the addition of specific compounds called ‘flame retardants’ (FR) in order to break the fire triangle. FR can be physically mixed with the polymer during the transformation process with the advantage of not modifying the chemical structure of the macromolecules. The FR used in this way constitutes the class of additive flame retardants. The FR function can also be introduced before (as monomers) or after polymerization (by chemical grafting) which requires a possibly consequent modification of the synthesis process but ensures no migration of the additive. The FR used in this way constitutes the class of reactive flame retardants. The diversity of polymers to protect, the different fire safety requirements according to the sectors of applications, the regulations and standards in effect require a wide variety of usable compounds. On the other hand, the optimization of the fire behavior often forces one to use two or more FR compounds combining different modes of action and developing synergistic effects to limit the quantities introduced and the deterioration of mechanical and functional properties of the material. The FR compounds can act either in the gas phase or in the condensed phase [2, 3].

22.2.2.1

Action in the Gas Phase

The FR compound upon decomposition generates active radicals in the flame zone trapping the hot radicals HO◦ and H◦ , responsible of its propagation and appreciably reducing the rate of the combustion reaction (flame poisoning). Another mode of action corresponds to the emission of inert gases such as H2 O, NH3 , CO2 by thermal decomposition of the FR compound. These gases then dilute the fuel in the flame zone causing the decrease of the rate of combustion of the material.

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22.2.2.2

Action in the Condensed Phase

Endothermic reactions taking place in the FR compound means that the temperature of the material remains sufficiently low, delaying the polymer decomposition and limiting the emission of combustible volatile gases. The FR compound can also be the site of reactions that promote the charring of the polymer and the formation of a carbonized layer more or less expanded to act as a barrier and protect the material from the flame, isolating the polymer from oxygen, reducing heat transfers and the emission of combustible volatile gases. In the so-called intumescent systems, a maximum expansion is seeked. The FR can also accelerate the decomposition and the rupture of the polymer chains, allowing the polymer to drip and thus to move away from the flame zone. In other cases, the FR compound participates only by the construction of the structure of the protective layer by mechanical reinforcement of the char. In this way, the formation of holes and cracks in the char favoring the emission of gaseous compounds is strongly limited.

22.3 Laboratory Fire Testing The role of fire tests is to characterize the burning behavior of the systems considered to improve the fire properties of polymer materials by simulating as best as possible the conditions of a real fire. The retained formulations must reach the conditions defined by the standards for each type of test. In the tests used in academic or industrial laboratories, the sample size may vary from a small strip of material whose dimensions do not exceed a few centimeters to large panels of about one meter or even beyond. They are used for either screening formulations during product development or testing manufactured products. Mainly, the fire tests should address the following questions: 1. Flammability: exposed to a flame or heat source, does the material ignite? What is the time of ignition? 2. Combustibility: once ignited, will it burn? What is the duration of combustion? 3. Heat: how fast will the fire spread? What will be the heat rate or power released? To determine this, one assesses the amount of heat released and the heat flow (rate of heat release), whose maximum value is the key indicator of the propensity of the material to propagate the flame. The flammability tests and standards having been already reviewed [2], only four of the most common testing methods used in the laboratory will be described hereafter. 22.3.1

Limiting Oxygen Index (LOI)

The value of the LOI index corresponds to the minimum percentage of oxygen in an O2 /N2 mixture to ensure a stable combustion with flame of a vertical sample during 3 minutes or on a 5 centimeters length under the conditions defined by the international standard: ISO 4589. According to this, the LOI is measured on 80 × 10 × 4 mm3 specimens placed vertically at the center of a glass chimney (Figure 22.3). The gaseous flow is admitted at the base of the chimney and homogenized by going through glass balls layers. After 30 seconds of purge, the upper end of the specimen is ignited using a burner. Materials showing an LOI less than 21 are said to be ‘combustible’, while those with an LOI higher than 21 are classified as ‘self-extinguishing’, meaning that the combustion can be sustained only with an external energy contribution. The flame retardant property increases with the value of LOI. Very simple to carry out, the LOI test is characterized by its high repeatability and reproducibility. Nevertheless, the user has to take into account that LOI values are measured at ambient temperature, and tend to decrease when the temperature increases. Consequently, this test will mainly provide information about the flame retardancy of samples which can be compared, and not indicate the exact behavior of a polymer in a real fire.

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specimen

mixture of gases (N2/O2)

Figure 22.3

22.3.2

Apparatus for LOI determination.

Epiradiator or ‘Drop Test’

The aim of this test is to submit the specimen to heat radiation, causing inflammation of the gas produced by the resulting degradation and/or whether there is a fall of burning drops. Square samples of 70 mm per side, so as to obtain a weight of over 2 g, are installed on a clearly-defined metal grid, and submitted to the radiation of a radiator located 3.0 cm above (Figure 22.4). On each ignition it is moved aside and replaced after extinction during the first 5 minutes; then, for a further 5 minutes, it remains in position. The measured parameters are: time to ignition, mean number of ignitions, combustion time; and the other determining elements are: the presence or not of burning drops, and the ignition of the cellulose wool placed under the test sample.

Radiator

30 mm

Radiant surface (500W) Sample sheet Metallic grid

135 mm

300 mm

Cotton

Figure 22.4

Epiradiator experimental set-up.

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127 specimen

9.5

305

burner

cotton

Figure 22.5

22.3.3

UL 94 V flammability test apparatus.

UL 94

Developed by Underwriters Laboratories (U.S. insurances laboratories), this test has been defined to serve as a preliminary indication of the flammability of plastic materials to be used as part of a device or appliance. UL 94 contains seven different tests and the most commonly used is UL 94V in which the ignitability and flame spread of a vertical specimen exposed to a small flame (50 W) is measured. The schematic experimental set up is given in Figure 22.5. A flame of 20 mm height is applied to the bottom of a sample with a length of 127 mm, width of 13 mm and thickness of 3 mm. After an application of 10 seconds, the flame is removed and the flame-out time is measured. The flame is then applied again for 10 seconds and removed and a second flame-out time is measured. If drops fall, the burner must be tilted through a maximum angle of 45◦ . The occurrence of burning drops causing the ignition of the cotton wool placed under the sample must be noted. Five different specimens should be tested. The test classifies materials as V0, V1 or V2. Mainly, for a V0 rating, burning stops within 10 seconds after each of the two flame applications and no flaming drips are allowed. For a V1 rating, burning stops within 60 seconds after each of the two flame applications and again no flaming drips are allowed. For a V2 rating, the same conditions of flame-out time than V1 are required but flaming drips are allowed. The complete UL standard can be obtained at http://ulstandardsinfonet.ul.com/. The international standard is: IEC 60695-11-10. 22.3.4

Cone Calorimeter

Cone calorimetry is one of the most effective medium-scale fire behavior tests for polymers in which the specimen is submitted to a given heat flux. It was developed by V. Brabauskas in the 1980s [4]. Figure 22.6 gives the scheme of the apparatus. The sample (100 × 100 × 4 mm3 ) is placed on a load cell in order to evaluate the evolution of mass loss during the experiment. A conical radiant electrical heater uniformly irradiates the sample from above. The combustion is triggered by an electric spark. The combustion gases pass through the heating cone and are captured by means of an exhaust dust system with a centrifugal fan and a hood.

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Laser photometer beam including temperature measurement Temperature and differential pressure measurements taken here Soot sample tube

Exhaust blower

Exhaust hood Gas samples taken here

Soot collection filter

Cone heater

Controlled flow rate

Spark igniter Specimen

Load cell

Figure 22.6

Cone calorimeter experimental set-up.

The principle of cone calorimetry is based upon the measurement of the decrease in oxygen content of combustion gases and the fact that the amount of heat released during combustion is proportional to the amount of oxygen consumed. For polymers (organic matter), the value of the proportionality constant is 13.1 kJ.g−1 consumed oxygen [5]. The measurement (using a paramagnetic analyzer) of the oxygen concentration and of the gas flow allows the calculation of the heat release rate (heat released per unit of time and surface area). Other measurements can be carried out simultaneously: time to ignition, time of flame out, mass loss rate, total heat released, amounts of emitted CO and CO2 , smoke density and total smoke released. The value of the maximum of the heat released rate curve is often taken into account to evaluate the fire properties as it characterizes the propagation rate of the flame and the spreading of fire. The cone calorimeter test is the subject of an international standard: ISO 5660.

22.4 Flame Retardant Additives 22.4.1

Hydrated Fillers

Metallic hydroxides such as aluminum tri-hydroxide (ATH) and magnesium di-hydroxide (MDH) decompose endothermically and release water. The endothermic decomposition of ATH (Al(OH)3 ) occurs between 180 and 200◦ C and leads to the release of water and the formation of alumina. MDH (Mg(OH)2 ) acts in the same way but its endothermic degradation occurs at a higher temperature (>300◦ C).

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These reactions have several effects on polymer combustion. They absorb energy (1050 kJ/kg in the case of ATH and 1300 kJ/kg with MDH), which leads to the cooling of the polymer; they form a thermally insulating protective coating (Al2 O3 and MgO); and the released water vapor dilutes combustible gases and forms a protective gas layer. They might also have a catalytic action on the degradation pathway of the polymer that produces a carbonized char residue. However, to achieve a high level of fire performance, it is necessary to use these hydrated fillers at high loading level. The flame retardant action of ATH and MDH is very effective up to 400◦ C. Beyond this temperature, the exothermic character of combustion predominates. Hydroxycarbonates can be also used as flame retardant additives owing to their endothermic decomposition due to the liberation of water and CO2 at high temperature. Hydromagnesite, for example, releases water and carbon dioxide on a larger temperature range [6] (200–550◦ C) than alumina trihydrate and magnesium dihydroxide (Associated heat = 800 kJ/kg). Zinc borates are used as flame retardants for polymers. Their endothermal decomposition, between 130 and 270◦ C, liberates water and boron oxide (B2 O3 ) that softens at 350◦ C and flows above 500◦ C leading to the formation of a protective vitreous layer. 22.4.2

Halogenated Flame Retardants

The thermal degradation of polymers releases reactive radical species such as H◦ and OH◦ , which maintain combustion by a cascade-chain mechanism in the gas phase. The halogenated flame retardants (RX; X= Br or Cl) are able to react with these reactive species (H◦ , R◦ and OH◦ ), stopping the chain decomposition, and then the combustion. Fluorinated organics are not used because they do not interfere with the polymer combustion process, they are generally more thermally stable than polymers and they cannot release halogen in the form of radical during the polymer combustion [7] whereas instead of when the iodinated organics degrade below the temperature of polymer degradation, bromine and chlorine can easily be released during the combustion process because of their weak binding energy with carbon. The most used halogenated flame retardants are tetrabromobisphenol A (TBBPA), polybromodiphenylethers (PBDE), hexabromocyclododecane (HBCD) and tetrabromophthalic anhydride (TBPA). The effective nonflammable species HX regenerated by the reaction of X◦ with the polymer (RH) have a physical action on the combustion mechanism by formation of a protective gaseous coating and by dilution of fuel gases. It catalyzes also the oxidation of the condensed phase that has a tendency to generate a protective solid char layer. R − Br → R◦ + Br◦ Br◦ + R H → R◦ + HBr HBr + H◦ → H2 + Br◦ HBr + OH◦ → H2 O + Br◦ Halogenated flame retardant can be incorporated in a blend with the polymer, used as polymerization monomers, copolymerized with virgin monomers or grafted on the polymer chain. When they are directly incorporated in the polymer chain, they can be used in a relatively low concentration that limits the damage caused by heterogeneous additives to the mechanical properties of the materials and reduces the problems of the migration of the flame retardant agent onto the material’s surface. 22.4.3

Phosphorus-based Flame Retardants

Phosphorus-based flame retardants are particularly effective in polymers containing oxygen or nitrogen, such as polyamide and polyesters [8, 9]. By thermal decomposition, phosphorus-based flame retardants generate

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phosphoric acid which condenses easily to give pyrophosphate structures that catalyze alcohol’s end-groups dehydration and leads to the formation of carbocations and carbon–carbon double bonds into polymer chains. The pyrophosphoric acids transform into metaphosphoric acid and into polyphosphates that take part, with the carbonized polymer remainders, in the formation at the surface of the material of a protective carbonized char layer. In the non-oxygen- or nitrogen-containing polymers that cannot be involved in the charring, a highly charring co-additive such as pentaerythritol [10] can be used. Phosphorus-based flame retardants can act also in the gas phase. They volatilize to form active radicals such as PO2 ◦ , PO◦ and HPO◦ , and act as scavengers of H◦ and OH◦ radicals. Ammonium polyphosphate (APP) is one of the most commonly-used phosphorus-based flame retardant products. It is a stable and nonvolatile inorganic salt of polyphosphoric acid and ammonia. Long chain APP starts to decompose above 300◦ C to generate polyphosphoric acid and ammonia. Short chain APP will begin to decompose above 150◦ C. It is then very important to adapt the APP chain length within the polymer decomposition temperature range. The free acidic hydroxyl groups, generated during APP thermal degradation, condense and lead to the formation of polyphosphoric acid and a crosslinked ultraphosphate structure [11]. The polyphosphoric acid react with the oxygen- or nitrogen-containing polymers and catalyzes their dehydration reaction leading to the char formation, that is combined with the generated phosphorus crosslinked structure to produce a protective char layer. Red phosphorus is a very efficient flame retardant agent in polymeric materials. In polymers containing oxygen and nitrogen, the thermal degradation of red phosphorus produces phosphoric anhydride or phosphoric acid and leads to the formation of polyphosphoric acid that catalyzes the polymer dehydration reaction and the char formation [12]. In non-oxygenated polymers such as polyethylene [13], red phosphorus acts in both the vapor and condensed phase [14, 15]. A fraction of red phosphorus depolymerizes to white phosphorus that volatilizes and forms active radicals such as PO◦ to scavenge H◦ and OH◦ species that quench radical processes. The other part of red phosphorus diffuses from the bulk to the burning surface where it is oxidized to phosphoric acid and its derivatives that act as a char-forming agent. Many organic phosphorus derivatives have flame-retardant properties; the number of those with commercial importance is limited by the processing temperature and the type of polymer. For example, the use of (alkyl-substituted) triaryl phosphates such as triphenyl phosphate or cresyl diphenyl phosphate is very limited in engineering plastics because of their high volatility. They may volatilize before the polymer thermal decomposition starts. Oligomeric phosphates, with lower volatility and higher thermal stability can be an alternative to limit the volatilization of these phosphorus derivatives during the thermal decomposition and permit their use for engineering plastics. The direct incorporation within the polymer chain structure by using reactive phosphorus flame retardants is also a solution against their migration towards the surface of the polymer and their volatilization during the thermal decomposition. Phosphorated melamine salts such as melamine phosphate and melamine pyrophosphate are also effective as polymer flame retardants. During heating, melamine phosphate dissociates in melamine, melamine pyrophosphate and melamine polyphosphate. At high temperature, melamine may decompose with ammonia elimination and the generation of thermally stable condensates, known as melam, melem and melon (Figure 22.7) [16] and the produced ammonia can dilute oxygen and combustibles in the gas phase. Melamine pyrophosphate and melamine polyphosphate react with the melamine condensation products to generate melam ultraphosphate and ammonium polyphosphate that dissociate above 300◦ C and release ammonia and free condensed hydroxyl groups that give a cross-linked structures (ultraphosphate). The melamine phosphate thermal decomposition process was established in detail by Levchik et al. [17].

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N N

NH2

N N

N

NH2 N N

NH2 2 N

N NH2

NH2 −NH3

NH2 N

N N

H2N

melamine

N

NH2 N NH melam

N N

−NH3 NH2

H 2N

N

N

N N

N

N

N

NH

−NH3 N

NH2 N

melem H2N

N N

N N

N

NH2

melon

Figure 22.7

Chemical structure of melamine and its derivatives (from Ref. [15]).

Some phosphorus-based flame retardants are also used in intumescent flame-retardant systems. The intumescent concept is based on the formation of an expanded protective carbonized layer on the surface of the burning material. It is well known that to generate intumescence, three components are required in the flame retardant system. The acid source such as ammonium polyphosphate promotes the dehydration of the carbonized agent (Pentaerythritol) leading to the char formation. The expansion is produced by the thermal decomposition of the blowing agent (melamine, guanidine, urea). The acid liberation must take place at a temperature below the decomposition temperature of the carbonizing agent and its dehydration must occur around the temperature of the polymer decomposition. The gas release must occur during the char formation to induce the expansion of the carbonized layer. The amount of charring produced during thermal decomposition strongly depends on the number of carbon atoms that contain the carbonized agent, while the number of its hydroxyl reactive sites determines the rate of the carbonized structure formation. 22.4.4

Nanometric Particles

The contribution to flame retardance of nanoparticles is different and depends on their chemical structures and geometries. Their use permits the reduction of the loading rate since the size reduction of the filler allows the interfacial area between the polymer and the filler to be largely increased. Nanoclays were the first nanoparticles used in polymer as flame-retardant agents. The organomodification of natural clays using organic cations (alkylammoniums and alkyl phosphoniums as well as alkyl imidazol(idin)iums cations) is required to promote their dispersion within polymers. The incorporation of a relatively low amount of (organomodified) nanoclay leads to the formation of a protective layer during the combustion. Upon heating, the viscosity and surface energy of the molten polymer/layered silicate nanocomposite decreases with increasing temperature and facilitates the migration of the clay nanolayers to the surface. Moreover, heat transfer promotes the thermal decomposition of the organomodifier and the creation of strongly protonic catalytic sites (Brønsted acid sites) on the clay layers, which catalyzes the formation of a stable char residue [18, 19]. The accumulation of the clay on the surface of the material therefore acts as a protective barrier that limits the heat transfer into the material, the volatilization of combustible degradation products and the diffusion of oxygen into the material.

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The dispersion state of the clay particles in the polymer matrix may have an important influence on the flammability [20]. The formation of an intercalated or exfoliated nanocomposite should be achieved to reach good flame-retardant properties. For example, the peak of heat release rate is decreased by 25% for EVA/5% R 30B nanocomposite in comparison with Na-MMT microcomposite and by 50% for the EVA/5% Cloisite the virgin polymer [21]. The formation of a barrier against heat and volatiles by clay nanolayer migration toward the material’s surface, followed by char formation but also by an increase of the melt viscosity in exfoliated nanocomposites are the main general fire-retardancy mechanisms in polymer/clay nanocomposites. They lead to the modification of the polymer nanocomposite fire properties, sometimes improving them and sometimes damaging them, depending on the type of fire tests applied and the nature of the measurements recorded. For example, in the cone calorimeter test, the incorporation of nanoclays retards mainly the peak of heat release rate in developing fires, but does not reduce the total heat evolved and often increases the flammability (time to ignition). The increased melt viscosity prevents dripping and promotes the formation of the char, but the char formed at the surface of the burning sample for UL-94 and LOI tests is not effective enough to stop the flame and the sample continues to burn slowly, leading to classification failure. Carbon nanotubes (CNTs) represent an interesting alternative to the use of conventional flame retardants and nanoclays. Thanks to their high aspect ratio, CNTs can percolate to form a network at very low loading. During combustion, carbon nanotubes promote the char formation by catalytic polymer dehydrogenation [22] that leads to the formation of a double bond on the polymer chain allowing polymer crosslinking. The chemical reactivity of CNT can be enhanced by the use of crushed MWNT [23] that contain radical species (or precursors) at their surface. The incorporation of 3% of such crushed carbon nanotubes, in ethylene-vinyl acetate (EVA) copolymer, leads to a large increase of time to ignition during the cone calorimeter test. The partial substitution of MWNT by crushed MWNT leads also to the same behavior. The quality of the char layer depends on the dispersion state of the CNTs and their loading rate. In PMMA [24], the formation of a compact and continuous char layer is obtained only for nanocomposites exhibiting good nanotube dispersion. On the other side, discontinuous residues were obtained for the nanocomposites containing poorly-dispersed nanotubes. In EVA [25], the use of HDPE-coated MWNT, obtained by the polymerization-filling technique [26], improves the dispersion state of CNT and leads to the formation of a more cohesive char during the cone calorimeter test. The formation of such an efficient and compact layer is favored also by the use of MWNT with high average aspect ratio (length-to-outer diameter ratio) [27]. Polyhedral Oligomeric SilSesquioxane (POSS) is recognized to be a precursor for the formation of thermally-stable ceramic materials. POSS is an inorganic silica-like nanocage ((RSiO1.5 )8 ) surrounded by eight organic groups (R) located at the corners. A large choice of organic groups is possible and allows POSS selectivity according to the chemical structure of the polymer matrix to be used. The incorporation of POSS in polymers affects the thermal stability and fire properties by the formation of a ceramic layer at the surface of the burning polymer. A char formation by catalytic dehydrogenation is also possible by using modified POSS containing metal atoms at one edge of the nanocage structure [28, 29]. In this case, metal-modified POSS plays the role of the metal dispersion agent and it was shown that finely dispersed metal-bearing POSS nanoparticles, at very tiny concentrations (1 wt%), strongly enhance PP char yield. Metallic oxide nanoparticles such as titanium oxide (TiO2 ), ferric oxide (Fe2 O3 ) and alumina (Al2 O3 ) also have a potential flame retardant action [30]. The incorporation of 5 wt% of nanometric TiO2 or Fe2 O3 leads to the improvement of the thermal stability of PMMA nanocomposites [31] and in the cone calorimeter test, the HRR values proved to depend on the filler content and decreased with the increase of nanoparticles content. The thermal stability and the flame-retardant improvement of the PMMA nanocomposites were attributed to different physical and chemical phenomena. The restriction of mobility of the polymer chains attributed to steric hindrance due to the presence of the nanoparticles and the adsorption of polymer on their surface via

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the methoxycarbonyl group can be at the origin of the modification of the polymer degradation pathway and the char formation.

22.5 Synergistic Effects of Fillers with Flame Retardant Additives 22.5.1

Definition of Synergistic Effects in Flame Retardant Systems

Synergistic effects are achieved when the performance level due to a mixture of additives xA + yB (x + y = 1) for a given property (P) is greater than that predicted for the linear combination (xPA + yPB) of the single effects of each additive (PA and PB). Conversely, antagonistic effects can be observed. Considering diverse flame-retardant properties such as limiting oxygen index or peak of heat release rate, for example, a synergistic effect can be achieved only for a specific property, and not for the others, since various phenomena can influence it. As a matter of fact, synergistic effects can arise not only from chemical interactions occurring in the gas phase, chemical interactions between species present in the gas phase and others present in the condensed phase, but also interactions between physical processes, such as limitations of mass or heat transfer and chemical ones. Hence, in multicomponent flame-retardant systems, and particularly if some components are polymers, it becomes hard to interpret the synergistic effects. However, some synergistic effects have been established and widely studied, such as brominated compounds/antimony oxide combinations, phosphorus/nitrogen compounds association.

22.5.2

Micronic Fillers and Flame Retardants

The high level of efficiency of halogenated compounds is counterbalanced by the correlated release of hydracids. In order to reduce the release of these corrosive gases, various combinations have been tested, particularly with compounds able to modify the flame chemistry. Antimony trioxide can be incorporated as a micronic filler in combination with halogenated compounds (used as additives or reactive flame retardants) or chlorinated polymers such as PVC. Sb2 O3 has very poor flame-retardant properties if used without halogenated species [32]. Nevertheless, a synergistic effect on auto-extinguishability can be noticed in relation to the formation of new species such as antimony oxyhalides, in the gas phase according to the following chemical reactions [33]: SbCl3 + H◦ → HCl + SbCl◦2 SbCl2 + H◦ → HCl + SbCl◦ SbCl + H◦ → HCl + Sb◦ These oxyhalides lead to the formation of antimony trihalides, which act as scavengers in the flame of very active radicals such as H◦ . In addition, Sb◦ radicals can also scavenge H◦ and OH◦ radicals. Sb◦ + OH◦ → SbOH SbOH + H◦ → SbO◦ + H2 SbO◦ + H◦ → SbOH

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For common polymers, specific Br/Sb ratios were highlighted to maximize the synergistic effect. Since a carcinogenic action of antimony trioxide has been suspected in addition to the restrictions on the use of halogenated compounds, the future of this combination is not very promising and the development of alternative FR systems has originated from the progressive phasing out of such FR combinations. The widespread use of halogenated compounds from the 1970s and the development of new phosphorus compounds able to play a flame-retardant role in polymers, has also led to investigations about possibilities of Br/P synergistic effects. Some mixtures of brominated and phosphorus compounds, or additives containing both phosphorus and bromine were developed in order to combine the action of bromine in the gas phase (as radical scavenger) and thus of phosphorus (e.g. phosphorus esters) in the condensed phase (as char layers promoter) [34]. Some synergistic effects between additive brominated flame retardants and hydrated mineral fillers have been investigated. The objective was to maintain the flame-retardant properties while reducing the release of corrosive gases in case of fire as well as to improve the mechanical properties of polymers filled with high percentages of mineral fillers. Montezin et al. [35] have combined decabromodiphenyl oxide/antimony trioxide blends with magnesium hydroxide. Synergistic effects were noticed on ignition time, moreover the UL 94 V0 classification was maintained as well as the LOI value in comparison with a composition containing only the brominated blend. As mentioned above, various forms of phosphorus compounds acting as flame retardants exist. Some of them are additives in a micronic form dispersed as fillers in polymers. They are able to produce synergisms either with nitrogen compounds or with metallic oxides or hydroxides. Concerning the associations between phosphorus and nitrogen compounds, synergisms have been noticed, for example between red phosphorus and melamine cyanurate in polyamides [36], also between melamine and triphenyl phosphate in poly(butylene terephthalate) [37]. Moreover, the action of intumescent systems is based on synergistic effects between their components containing nitrogen and phosphorus such as ammonium polyphosphate, melamine or melamine salts, such as melamine phosphate, pyrophosphate or cyanurate. The selection of an adequate association between nitrogen and phosphorus compounds would often depend on the nature of the polymer, whose decomposition pathway is modified by the action of these components, particularly in the condensed phase. The formation of nitrogenphosphorus intermediates can promote the production of phosphoric acid and then the phosphorylation of the polymer [38]. Due to the more reactive character of P-N bonds than P-O ones towards the phosphorylation process, the phosphorus–nitrogen combination enhances the phosphorus content in the condensed phase, allowing the promotion of char through the formation of crosslinked structures. Several fillers have been incorporated in order to produce synergistic effects in intumescent systems containing ammonium polyphosphate. The flame-retardant effect of only such fillers in the polymer is poor or nonexistent. Le Bras et al. [39] have highlighted synergistic effects on the LOI of polystyrene, by combining a zeolite 13X, usually employed as catalyst or molecular sieve, with an intumescent system made up of ammonium polyphosphate (APP) and pentaerythritol. Recently, Basfar et al. [40] have shown the interest of the incorporation of talc in intumescent compositions containing APP in Ethylene Vinyl Acetate. Other types of synergism involve blends of metallic hydroxides acting as hydrated fillers with other kinds of fillers. For example, in order to reduce the global amount of fillers, the endothermic effect of hydroxides can be improved by using zinc borates which have often been used as synergistic agents in polyolefins. Magnesium hydroxide has been partially replaced by zinc borate in EVA by Carpentier et al. [41] and synergistic effects have been found on peaks of HRR and LOI values. It has been shown that the decomposition of magnesium hydroxide could catalyze the decomposition of zinc borate leading to boron oxide at lower temperatures. A vitreous and cohesive layer made up of MgO and boron oxide was then formed, ensuring a stronger physical barrier effect than MgO alone. Durin-France et al. [42] have shown the interest of the addition of lamellar talc particles in mineral ternary blends containing also magnesium hydroxide and zinc borate in EVA.

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The development of the use of fillers having nanometric dimensions in polymers has not only entailed investigations about the interest of the combination of these new classes of fillers with flame retardant in additive combinations, but also to the creation of complex compositions in which the nanometric fillers are a part of multicomponent flame-retardant systems. 22.5.3

Nanometric Fillers and Flame Retardants

In the 1990s, various research groups highlighted the flame-retardant properties conferred by nanoparticles and mainly by organomodified lamellar silicates (OMLS) [43–45]. Nanoparticles incorporated at percentages up to 10 % wt reduced polymer flammability and enhanced the formation of char. The mechanisms of fire reaction involving nanoparticles are rather different according to their nature and treatments to those exhibited by usual flame retardants and correspond to various physical, physico-chemical or chemical processes:

r r r r r

formation of mass transfer barriers for combustibles volatiles and oxygen by migration of nanoparticles towards the surface exposed to heat flux formation of insulating char structures (possibly expanded) through catalytic effects, able to limit the temperature, or able to dissipate the incident heat by radiative emission modification of heat diffusivity through the material restriction of macromolecular mobility and increase of viscosity trapping of radicals released from the thermal degradation and combustion.

However, the incorporation of nanoparticles in polymers by various processes is unable to meet fire performance standards in comparison with usual flame retardants (FR), such as hydrated minerals, halogen, phosphorus or nitrogen compounds. In consequence, the concept of combining nanoparticles with flame retardants, mainly non-halogenated, has emerged, in order to generate synergies, leading also to the possible reduction of the global loading in polymers. This last objective aimed particularly at combining high levels of flame retardancy, compatible with demanding standards, the conservation and even the improvement of some mechanical properties, mainly stiffness and reinforcement, due to the presence of nanoparticles. Surface modifications of nanoparticles aim to promote a good dispersion of the nanoparticles into the polymer matrix (intercalated or exfoliated morphologies for surface modified layered silicates), also in the presence of usual flame retardants allowing strong synergistic effects to be achieved. Moreover, some interactions between nanoparticles and recognized flame retardant can occur, particularly reactions leading to new mineral species, depending on the flammability conditions. The development of original strategies for surface modifications of nanoparticles with compounds having a flame-retardant activity could provide a new field of research on flame-retardant systems. The use of original phosphorus, nitrogen or halogen containing modifiers instead of alkylammonium ions for layered silicates seems promising. A flame-retardant action conferred by the surface modifier can be combined to an action on composite morphology, particularly for a polymer blend as host polymer. In addition, some FR compounds currently used as micronic fillers can present an enhanced reactivity at submicronic scale and possibly with surface treatments. For example, new varieties of metallic hydroxides were synthetized, either able to play a similar role as some lamellar nanoparticules or to act as conventional hydrated FR fillers. Finally, the use of flame-retardant systems in which organomodified nanoparticles, exhibiting various flameretardant mechanisms, are associated, could also represent new alternatives to conventional flame-retardant systems.

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The association of natural layered silicates with intumescent flame-retardant systems represents one of the main contributions of the combined use of nanoparticles and flame retardants. Moreover, combinations of layered silicates with other phosphorus compounds have been studied and have led to significant improvements for fire reaction.

22.5.3.1

Combination with Usual Flame Retardants

The main category of nanoparticles involved in combinations with usual phosphorus compounds, metallic hydroxides, and even with halogenated compounds [46] are organomodified silicates (OMMT). The flame-retardant properties of OMMT have been studied for a wide range of polymers, particularly in thermoplastics. Depending on the nature of the polymer, the decomposition pathway and decomposition products may change [19] with the formation or the enhancement of charred structure, caused by crosslinking processes possibly catalyzed by the nanoparticles. The main categories of clays studied are cationic clays and mainly 2:1 silicates belonging to the category of dioctahedral smectites: montmorillonite or bentonite [47]. From purified materials, cation exchange procedures using mainly alkylammonium ions can lead to organophilic materials able to disperse into polymers. Only resulting exfoliated or intercalated structures [48] can allow high levels of flame retardancy to be achieved [43–45]. The combined use of OMMT with components of intumescent compositions such as ammonium polyphosphate was widely reported for various host polymers or blends, in which the role of the carbon source is played by a polymer (PA6, EVA, thermoplastic polyurethanes). Owing to its ability to form a stable char, EVA has been investigated as a polymer matrix for both ammonium polyphosphate and organomodified silicate. Due to the ability to achieve a nanocomposite structure in PA6 showing advantageous mechanical properties, the direct incorporation of PA6 [49] containing nanoparticles at a percentage of 2 %wt into EVA containing APP (at a global filler loading of 60 %wt) allowed LOI and HRR values to be strongly improved. Intumescent flame retardant-montmorillonite compositions for PP were also investigated by Tang et al. [50]. An intumescent flame-retardant system based on APP, pentaerythritol (PER) and MPP (melamine pyrophosphate) was incorporated in a PP containing montmorillonite and a compatibilizer (hexadecyltrimethyl ammonium bromide). Synergistic effects on HRR values between clay and intumescent FR (IFR) system were observed (Figure 22.8). OMMT have been associated with other phosphorus compounds such as organic phosphates and red phosphorus. Investigations about synergistic effects of hydrotalcite, a lamellar double hydroxide, with microencapsulated red P have been carried out by Du et al. [51] in EVA. The percentage of hydrotalcite was comprised between 25 and 38 %wt, while the percentage of phosphorus varied between 2 and 15 %wt, for a global loading at 40 %wt. Increasing amounts of red P up to 10% led to a decrease of HRR values and to an increase for the flame-out time at cone calorimeter tests (Figure 22.9). Aromatic phosphates are widely used in engineering plastics such as styrenics and polyesters. Polystyrene–clay nanocomposites combined with various phosphorus-containing FR were prepared by Chigwada et al. [52]. The clay modifier was dimethyl benzyl hydrogenated tallow ammonium and the OMMT was found intercalated in the polymer. Cone calorimeter measurements evidenced synergistic effects between phosphate flame retardants and organomodified clay. It was also shown that a percentage of 30% for the phosphate is needed to achieve a V0 rating at UL94V. In order to limit the global amount of metal hydroxides incorporated in usual flame-retarded polymers, particularly polyolefins, combinations between these flame retardants and OMMT aimed to reduce the global loading at constant fire behavior, particularly for EVA or LDPE. Barrier effects produced by the clays to

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Figure 22.8 HRR for PP/OMMT nanocomposites with IFR. Reprinted from [50]. Copyright (2003) with permission from John Wiley & Sons.

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Figure 22.9 Effect of red phosphorus (R) and hydrotalcite (H) percentage on HRRs of EVA/hydrotalcite/red phosphorus composites. Reprinted from [51]. Copyright (2006) with permission from Elsevier.

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regulate the water release and cohesion of the oxide residue formed at the surface of sample after polymer ablation have been expected to promote synergistic effects. Several authors [53–56] have presented synergistic effects on fire properties between organomodified layered silicates and metal hydroxides (magnesium hydroxide (MH) or alumina trihydrate (ATH)) in polyolefins. The quality of char formed for metal hydroxide/MMT combinations seems to act strongly on the efficiency of the barrier effect. Incorporation of silica with MH and OMMT by partial substitution of OMMT, was studied by Ferry et al. [53], and Laoutid et al. [57] to improve the cohesion of charred and expanded structures. Despite the presence of cracks generated by silica in the char, fire behavior as studied by cone calorimeter has been improved. Among the categories of nanoparticles being incorporated in FR systems, carbon nanotubes play a particular role due to their high aspect ratio and their ability to percolate at very low percentage to form a network in host polymers. Regardless of the nature of these nanotubes (single or multi-wall), their dispersion state in the host polymer is crucial as shown by Kashiwagi et al. in PMMA [24] and its improvement is a challenge to achieve the best fire performance of the corresponding nanocomposites. The flame-retardant properties of polymer/CNTs nanocomposites appear to be governed by antagonistic processes, since the CNT network acts as a shield and re-emits much of the incident radiation back into the gas phase, and the presence of carbon nanotubes increases the polymer thermal conductivity of the materials. CNT surface reactivity also leads to some modifications and improvement of nanocomposites’ fire properties, especially in the case of crushed MWNTs [58]. Combinations of CNT with organomodified clays and with ATH in EVA cable compositions, have been carried out by Beyer et al. [59–61]. A synergistic effect for flame retardancy between CNTs and organoclays was observed, also with a reduction of the crack density. It was concluded that organoclays play an active role in the formation of a compact char, reinforced by the CNTs due to their high aspect ratio. The synergistic effects of CNTs combined with organoclay were also kept from compositions containing ATH. Oxide nanoparticles have been combined with flame retardants, and mainly with phosphorus, and in ternary compositions with metal hydroxides and organomodified layered silicates. The interest of silica and oxide nanoparticles is particularly due to their effects on viscosity at molten state, and reactivity during the degradation stages of polymers. Pyrogenated silica was combined by Fu et al. [62] with MDH in EVA. The partial replacement of MH (at constant loading of 60 wt%) by a given amount of fumed silica increased the LOI value and allowed to maintain the V-0 rating in the UL-94 test. Laachachi et al. [63, 64] have combined oxide nanoparticles (Fe2 O3, Al2 O3 and TiO2 ) and organoclays to improve thermal stability and fire retardancy of PMMA by melt blending. A synergistic effect was also found by the combination of TiO2 and organoclays resulting mainly in an increase of the ignition time and the reinforcement of the barrier effect of organomodified clays. These results were ascribed to the promotion of charring due to OMMT together with thermal properties and high specific surface area of TiO2 , leading to a modification of heat transfer. The same authors have also successively investigated the combinations between metal oxide nanoparticles (Al2 O3 and TiO2 ), respectively with APP (or APP + melamine phosphate additive) and then with phosphinate derivatives in PMMA [64, 65] (Figure 22.10).

22.5.3.2

New Innovative Flame-Retardant Systems Based on Nanoparticles

The incorporation of nanoparticles in flame retarded polymers can lead to innovative composite materials when microstructures, interfaces and chemical reactivity can be controlled either by the surface properties or reactivity of nanoparticles, or by the use of interfacial agents or nanoparticle modifiers. Some of these interfacial agents or modifiers are often oligomers or macromolecules, with structures allowing flame retardancy (presence of phosphorus, nitrogen or silicon) as well as structural properties to be improved, since reactivity

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or compatibility with the host polymer(s) are optimized. Synergistic effects can arise only from the use of these interfacial components, at very low percentages of incorporation. The combination of an intumescent flame retardant (IFR) system with OMMT and a PA6/EVA alloy in a PP containing maleic anhydride polypropylene (MAPP) has been investigated [66, 67]. The improvement of flame retardancy and mechanical properties of the PP compositions containing only IFR and modified clay was achieved using the compatibilizing effect of the EVA/PA6 alloy. A copolymer was formed from the reaction between MAPP and PA6. Synergistic effects with the IFR system were observed. Other compatibilizing agents other than MAPP were studied in IFR systems for PP containing ammonium modified OMMT as synergist and PA6 as charring agent. Ma et al. [68] used carboxylated polypropylene (EPP) as a reactive compatibilizer for PP/PA6 blend, leading to co-crystallization processes. The presence of OMMT with an optimum value of 4 %wt, allowed dripping behavior at flammability tests to be prevented. In flame-retardant compositions for polypropylene, Marosi et al. [69, 70] have incorporated reactive blends of several OMMT with a cationic surfactant, a silicone, a polyborosiloxane elastomer (BSil) and an IFR system made of APP and polyol in PP. Microthermal analysis and rheological measurements proved the existence of a coating on clays by the elastomer. Raman spectroscopy showed that thermal degradation of filled PP samples resulted in accumulation of clay particles at the surface when coated by BSil elastomer. Moreover, the cone calorimeter residue with better BSil was found to be more flexible and durable. This improvement was ascribed to the thermal stability of silicone in comparison with the cationic surfactant. It was also concluded that, during thermal degradation, the particle surface treatment could act as a stable and protective layer, and a carrier delivering the nanoparticles to the surface. The development of nanometric hydroxides (ATH and MDH) and synthetic lamellar silicates such as lamellar double hydroxides (LDH) has led to new possibilities of surface or interfacial modifications. Intercalation of LDH by phosphorus containing ions was performed in EVA by Ye and co-workers [71] to produce a MgAl-PO4 hydrotalcite (HTP). Percentages of LDH were varied between 40 and 60 %wt and comparisons were made with MgAl-CO3 hydrotalcite. LOI values of MgAl-PO4 composite were found to be 2% higher than those of a reference MgAl-CO3 hydrotalcite, able to release carbon dioxide. Moreover, V1 rating for UL94V test could be achieved. FTIR spectra revealed that in MgAl-PO4 composition,

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Figure 22.11 Heat release rate curves obtained in cone calorimeter tests for formulations #2, 5, 8, 9 (respectively: PET/PC (80:20), PET/PC (80:20) + 5% MMT-P, PET/PC (80:20) + 4% TPPi, PET/PC (80:20) + 5% MMT-P+ 4% TPPi). Reprinted from [73, 74]. Copyright (2008) American Chemical Society.

the charred structure was more compact, involving P-O-P and P-O-C bonds and had faster formation than of MgAl-CO3 . In order to reduce the volatile character of some phosphorus compounds, an innovative strategy consists of intercalating them between the platelets of a layered silicate. Kim et al. [72] have intercalated triphenyl phosphate (TPP) in a montmorillonite incorporated in ABS. It was found that TPP intercalated in the clay presents a higher evaporation temperature compared to TPP. In addition, thermal stability of ABS was increased by incorporation of TPP-modified clay. The incorporation of epoxy resin at global loading of all components, kept constant at 15 %wt, resulted in a significant increase of LOI. It was suggested that this enhancement could be related to the compact character of chars formed. A phosphonium-modified montmorillonite (MMT) or a phosphorus-functionalized submicronic kaolin, have been incorporated in PET or PC-PET blend [73, 74]. The phosphorus compounds used were triphenyl phosphite (TPPi) or its methyl triphenoxy phosphonium iodine salt. Surface hydroxyls of kaolinite reacted with the first one, while the ionic form was exchanged with sodium ions of the MMT. The combination in an 80/20 wt% PET/PC blend of 4 %wt P-modified MMT and 5 %wt TPPi present as additive was advantageous in comparison with the same components used separately (Figure 22.11). This combination allowed V0 instead of V2 rating at UL94V test to be achieved. These results were ascribed to the compatibilizing effect of modified OMMT as well as to the chain extender action of TPPi. A last route allowing phosphorus having an increased flame-retardant effect into host polymers filled with organomodified layered silicates is to perform in situ polymerization using phosphorated monomers. For example, a novel phosphorus-containing PET copolymer was synthesized by Ge et al. [75] by in situ intercalation polycondensation of terephtalic acid, ethylene glycol, and 2-carboxyethyl(phenylphosphinic) acid (HPPPA) with a montmorillonite. It has been shown that the nanocomposite had a better flame retardancy than neat PET-co-HPPPA. An increase of LOI values (from 31.4 to 34.0) was noticed for a very low montmorillonite content. For a clay loading of 2 %wt, V0 rating could be achieved. Another synthesis using a diol grafted by a DOPO structure (DDP) as monomer instead of HPPPA, also led to interesting results.

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22.6 Conclusion The flammability of plastic materials is a very complex scientific and technological problem for which no single solution can be found, particularly regarding the wide range of polymer and composite materials and multiphase polymer systems used. Many flame-retardant systems acting by diverse mechanisms have either already been studied or are currently under development to ensure greater safety in a world where fire prevention remains a major issue and there is toxicological as well as environmental concern about materials and chemical compounds. Hence, flame-retardant systems containing halogenated compounds are progressively being phased out due to health and environmental concerns. Other systems containing relatively low amounts of nanoparticles and flame retardants from various families show very promising results. Among non-halogenated fire retardant additives, phosphorus and nitrogen-based compounds have proved to be very powerful solutions, especially in matrices containing oxygen or nitrogen atoms in their backbone, while silicon-based compounds also appear to provide efficient solutions. Moreover, with the emergence of new techniques such as cone calorimetry, research into fire retardancy is becoming a less and less empirical science, and a much better understanding of the fire reaction of multiphase polymer systems is arising from the scientific studies carried out, allowing a better development or selection of the most relevant flame-retardant systems, which are becoming more and more complex. This complexity is frequently driven by the research of synergistic effects, allowing to maximize the flame retardant as well as mechanical or other properties, and to minimize the percentage of fillers and additives in the polymer. The most interesting synergistic effects have been observed for multiphase systems containing nanoparticles. Since the majority of research works carried out on flame retardancy on nanocomposites have dealt with organomodified layered silicates, the most investigated combinations have concerned this class of nanoparticles with usual flame retardants in usual polymers, particularly polyolefins. Hence, the development of intumescent FR systems has led to the development of compositions combining organomodified lamellar silicates and intumescent systems with ammonium polyphosphate and co-synergists able to promote the formation of a stable and expanded char layer. Other families of nanoparticles such as carbon nanostructures (particularly carbon nanotubes) and nanometric oxide particles offer promising solutions. Due to their high effectiveness on thermal stability and fire reaction, carbon nanotubes have been associated mainly with metallic hydroxides, and also with organomodified lamellar silicates. The interest of nanooxides lies in the large range of possible surface modifications, according to the nature of host polymers, as well as their reactivity towards phosphorus flame retardants such as phosphates or phosphinates, leading to new chemical structures able to promote and reinforce char layers formed at the polymer surface. The future of flame retardant for multiphase polymer systems will certainly integrate an increasing role for the control of microstructures, interfaces and chemical reactivity of components. This will be performed by innovative surface modifications of micronic or nanometric fillers, and by the use of compatibilizing agents for the polymers also present in such systems.

Acknowledgements F. Laoutid would like to thanks R´egion Wallonne and the European Union (FEDER, FSE) for financial support in the frame of “Fonds structurels europ´eens 2007–2013 – FEDER Convergence” and “EU Seventh Framework Programme FP7”.

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References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31.

32. 33. 34. 35. 36. 37. 38.

E. Stauffer, Science and Justice, 43(2003) 29–40. J. Troitzsch, International Plastics Flammability Handbook, 2nd Edition, Hanser Publishers, Munich (1990). A. R. Horrocks, D. Price, Fire Retardant Materials, CRC Press, Boston (2001). V. Brabauskas, NBSIR 82–2611, US. Natl. Bur. Stand. (1982). C. Huggett, Fire Mater. 4(1980) 61. L. Haurie, A. I. Fernandez, J. I. Velasco, J. M. Chimenos, J. R. Tico-Grau, F. Espiell, Macromol. Symp., 221(2005) 165–174. S. Levchik, Introduction to flame retardancy and polymer flammability, in Flame Retardant Polymer Nanocomposites, A. B. Morgan and C. A. Wilkie (Eds) Wiley, Hoboken (2007). E. D. Weil, Encyclopedia of Polymer Science and Technology. Wiley Interscience, New York, Vol.11 (1986). A. M. Aronson, Phosphorous Chemistry, ACS Symposium, 486,(1992) 218. G. Camino, L. Costa, L. Trossarelli, Polym. Degrad. Stab. 6(1984) 243–252. G. Camino, N. Grassie, I. C. McNeill, J. Polym. Sci.: Polym. Chem. Ed., 16(1978) 95–103. A. Granzow and J. F. Cannelongo, Journal of Applied Polymer Science, 20(1976) 689–701. E. N. Peters, J. Appl. Polym. Sci., 24(1979) 1457–1464. U. Braun, B. Schartel, Fire retardancy mechanisms of red phosphorus in thermoplastics, Proceeding of the Additives 2003 Conference, San Francisco, CA (2003). A. Ballistreri, G. Montaudo, C. Puglisi, D. Vitalini, J. Polym. Sci., Polym. Chem. Ed., 21(1983) 679–688. L. Costa, G. Camino, D. Luda, P. Cortemiglia, Fire and Polymers, Chap. 15, ACS Symposium Series 425, Washington DC (1990). S. V. Levchik, L. Costa and G. Camino, Polym. Degrad. Stab. 43(1994) 43–54. R. Song, Z. Wang, X. Meng, B. Zhang, T. Tang, J. Appl. Polym. Sci. 106(2007) 3488–3494. B. N. Jang, M. Costache, C.A. Wilkie, Polymer, 46(2005) pp.10678–10687. H. O. Pastore, A. Franche, E. Boccaleri, L. Marchese, G. Camino, Macromol. Mater. Eng. 289(2004) 783–791. S. Bourbigot, S. Duquesne, C. Jama, Macromol. Symp. 233(2006) 180–190. A. Fina, S. Bocchini and G. Camino, Polym. Degrad. Stab, 93(2008), 1647–1655. S. Peeterbroeck, F. Laoutid, B. Swoboda, J.-M. Lopez-Cuesta, N. Moreau, J.B. Nagy, M. Alexandre, Ph. Dubois, Macromol. Rapid. Commun. 28(2007) 260–264. T. Kashiwagi, F. Du, K. I. Winey, K. M. Groth, J. R. Shields, S. P. Bellayer, H. Kim, J. F. Douglas, Polymer 46(2005) 471–481. S. Peeterbroeck, F. Laoutid, J.-M. Taulemesse, F. Monteverde, J.-M. Lopez-Cuesta, J. B. Nagy, M. Alexandre, Ph. Dubois, Adv. Funct. Mater., 17(2007) 2787–2791. D. Bonduel, M. Mainil, M. Alexandre, F. Monteverde, Ph. Dubois, Chem. Commun, (2005) 781–790. B. H. Cipiriano, T. Kashiwagi, S. R. Raghavan, Y. Yang, E. A. Grulke, K. Yamamoto, J. R. Shields, J. F. Douglas, Polymer, 48(2007) 6086–6096. A. Antonov, M. Yablokova, L. Costa, A. Balavanovich, G. Levchik, S. Levchik, Mol. Cryst. Liq. Cryst. Sci. Technol. Sect. A 353(2000) 203–212. A. Fina, D. Tabuani, A. Frache, G. Camino, Polymer, 46(2005) 7855–7866. A. Laachachi, M. Cochez, M. Ferriol, J.M. Lopez-Cuesta, E. Leroy, Mater. Lett. 59(2005) 36–39. A. Laachachi, Ph.D thesis, Polym´ethacrylate de m´ethyle (PMMA) : D´eveloppement de nouveaux syst`emes retardateurs de flamme a` base de nanocharges min´erales. Recherche de synergies avec des montmorillonites et des compos´es phosphor´es, Universit´e de Metz, France, (2005). A. Laachachi, M. Cochez, M. Ferriol, E. Leroy, J. M. Lopez Cuesta, N. Oget Polym. Deg. Stab., 85(2004), 641–646. J.W. Lyons, The Chemistry and Uses of Flame Retardants, R.E. Krieger Pub. Comp. (1997). J. Green, Polym. Deg. Stab., 54(1996) 189–193. F. Montezin, J.M. Lopez-Cuesta, A. Crespy, P. Georlette, Fire and Mater., 21(1997) 245–252. Y. Liu, Q. Wang, Polym. Deg. Stab. 91(2006) 3103–3109. J. Xiao, Y. Hu, L. Yang, Y. Cai, L. Song, Z. Chen, W. Fan, Polym. Deg. Stab., 91,2006, 2093–2100. H. Horacek, R. Grabner, Polym. Deg. Stab., 54(1996) 205–215.

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39. M. Le Bras, S. Bourbigot, Fire retarded intumescent thermoplastics formulations, synergy and synergistic agents, a review, in Fire Retardancy of Polymers, The Use of Intumescence, Royal Society of Chemistry, 1998. 40. A. A. Basfar, J. Mosnacek, T. M. Shukri, M. A. Bahattab, P. Noireaux, A. Coudreuse, J. Appl. Pol. Sci., 107(2008) 642–649. 41. F. Carpentier, S. Bourbigot, M. Le Bras, S. Delobel, Polym. Int. 49(2000). 1216–1221. 42. A Durin-France, L Ferry, J-M Lopez Cuesta, A Crespy, Polym. Int. 49,(2000) 1101–1105. 43. E.P. Giannelis Polymer layered nano-composites, Adv. Mater., 8,(1996) 29–39. 44. J. Gilman, T. Kashiwagi, S. Lomakin, E. Giannelis, E. Manias, J. Lichtenhan, P. Jones, in Fire Retardance of Polymers: The Use of Intumescence, Royal Society of Chemistry, London, 1998. 45. J. Gilman, T. Kashiwagi, J. Lichtenhan, Nano-composites: a revolutionary new flame retardant approach, Sampe J., 33(1997) 40–46. 46. Y. Hu, S.F. Wang, Z.H Lin, Y.L. Zhuang, Z.Y. Chen, W.C. Fan, Macromol. Mater. Eng. 288(2003) 272–276. 47. F. Bergaya, B.K.G. Theng, G. Lakaly (Eds) Handbook of Clay Science (2006) Elsevier. 48. R.A. Vaia, H.Ishii, E.P. Giannelis, Chem.Mater., 5(1993) 1694–1696. 49. S. Bourbigot, M. Le Bras, F. Dabrowski, J.W. Gilman, T. Kashiwagi, Fire Mater. 24(2000) 201–208. 50. Y. Tang, Y. Hu, Y.S. Wang, Z. Gui, Z. Chen, W. Fan, Polym. Int., 52(2003) 1396–1400. 51. L. Du, B. Qu, Z. Xu, Polym. Deg. Stab., 91(2006) 995–1001. 52. G.Chigwada, C.A. Wilkie, Polym. Deg. Stab., 80(2003) 551–557. 53. L. Ferry, P. Gaudon, E. Leroy, JM Lopez Cuesta, Intumescence in EVA copolymer filled with magnesium hydroxide and organoclays, in Fire Retardancy of Polymers, Royal Society of Chemistry, London, 2005. 54. J.Zhang, C.A. Wilkie, Polym. Adv. Tech. 16,(2005) 549–553. 55. N. Ristolainen, U. Hippi, J. Seppala, A. Nikanen, J. Ruokolainen, Polym. Eng. Sci., 45(2005) 1568−1575. 56. G. Beyer, Fire Mater., 25(2001), 193–197. 57. F. Laoutid, L. Ferry, E. Leroy, J.M. Lopez-Cuesta, Polym. Deg. Stab. 91(2006), 2140–2145. 58. S. Peeterbroeck, F. Laoutid, B. Swoboda, J.-M. Lopez-Cuesta, N. Moreau, J.B. Nagy, M. Alexandre, Ph. Dubois, Macromol. Rapid. Commun., 28(2007) 260–271. 59. G. Beyer, Polym. Adv. Technol. 17(2005) 218–225. 60. G. Beyer, Fire Mater. 29(2005), 61–69. 61. F. Gao, G. Beyer, Q. Huan, Polym. Deg. Stab. 89(2005) pp. 559–564. 62. M.Fu, B.Qu, Polym. Deg. Stab. 85(2004) 633–639. 63. A. Laachachi, E. Leroy, M. Cochez, M. Ferriol, J.M. Lopez-Cuesta, Polym. Deg. Stab. 89(2005) 344–352. 64. A. Laachachi, M. Cochez, E. Leroy, P. Gaudon, M. Ferriol, J. M. Lopez Cuesta, Polym. Adv. Technol. 17(2006), 327–334. 65. A. Laachachi, M. Cochez, E. Leroy, M. Ferriol, J.M. Lopez Cuesta, Polym. Deg. Stab. 92(2007) 61–69. 66. Y. Tang, Y. Hu , B. Li, L. Liu, Z. Wang, Z. Chen, W. Fan, J. Polym. Sci. A Polym. Chem. 42(2004) 6161–6173. 67. Y. Tang, Y. Hu, L. Song, R. Zong, Z. Gui, W. Fan, Polym. Deg. Stab, 91(2006) 234–41. 68. Z.L. Ma, W.Y. Zhang, X.Y. Liu, J. Appl. Polym. Sci., 101(2006) 739–746. 69. G. Marosi, A. Marton, A. Szep, I. Czontos, S. Keszei, E. Zimonyi, A. Toth, X. Almeras, M. Le Bras, Polym. Deg. Stab., 82(2003) 379–385. 70. S. Keszei, Sz. Matko, Gy. Bertalan, P. Anna, Gy. Marosi, A. Toth, Eur. Polym. J., 41(2005) 697–705. 71. L. Ye, B. Qu, Polym. Deg. Stab. 93(2008) 918924. 72. J. Kim, K. Lee, J. Bae, J.Yang, S. Hong, Polym. Deg. Stab 79(2003) 201–207. 73. B. Swoboda, E. Leroy, L. Ferry, N. Kerboua, J.M. Lopez Cuesta, Fire properties of alloys and composites based on recycled polyethylene terephtalate, Proceedings of the 19th BCC Conference on Flame Retardancy, M. Lewin (Ed.), Business Communications Co Editions, Norwalk, CT, USA (2008). 74. B. Swoboda, E. Leroy, F. Laoutid, J.M. Lopez Cuesta, Flame retardant PET/PC blends compatibilized by organomodified montmorillonites, Proceedings of the ACS Conference, New Orleans (2008). 75. X.Ge, D.Wang, C. Wang, M.Qu, J.Wang, C. Zhao, X.Jing, Y.Wang, Eur. Polym. J. 43(2007) 2882–2890.

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23 Applications of Selected Multiphase Systems Igor Nov´ak Polymer Institute, Slovak Academy of Sciences, Bratislava, Slovakia

Volkan Cecen Department of Mechanical Engineering, Dokuz Eylul University, Bornova, Izmir, Turkey

Vladim´ır Poll´ak Polymer Institute, Slovak Academy of Sciences, Bratislava, Slovakia

23.1 Introduction Multiphase systems present combinations of two or more organic and/or inorganic materials. The special case of multiphase systems presents polymeric composites. In these systems, one of the materials is called the reinforcing phase, obviously in the form of organic or inorganic filler, i.e. fibers, sheets, or particles, which is embedded in the polymer matrix phase. The polymeric composites combine the strength of the reinforcement filler material and the toughness of the polymeric matrix. The strength of the polymeric composites is determined by the amount, arrangement, and type of filler concentration in the polymeric matrix. Over the past 40 years the applications of polymeric composites have increased, and at present the composite materials have various applications. Polymeric composites are widely used as materials in the production of many industrial products, e.g. for car construction in the automotive industry, and in marine, aircraft and space applications, where a wide variety of reinforced polymeric materials have begun to replace legacy materials.The commercial application of structural composites has an extensive history in the marine, automotive, sporting goods, and construction industries. The energy sector continues to be directed by new applications of composite materials to achieve unprecedented performance gains. Many polymeric composites are used for medical applications, for example in clinical practice to restore teeth. In most applications, dental composite consists of a polymeric acrylic or methacrylic matrix. Most human tissues, such as bones and, skin, are based on polymeric composites developed from constituents

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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whose amount, distribution, and properties determine the behavior of the material. Polymeric composites are used for the production of prostheses able to mimic biological tissues, restoring the mechanical function of the original tissue or organ. Composites are proving advantageous as components of electronics manufacturing systems, and as a commodity frequently used in the electronic industry, e.g. for production of printed circuit boards. Composites are now a commodity which dominates printed circuit board substrates and the production of parts of personal computers. Epoxies, polyurethanes, and polyamides are used for the encapsulation of many electronic components. Different cases of polymeric composites applications in various fields of human activities, i.e. in automotive, marine, aircraft, aerospace, human medicine, electrical engineering, and electronics, are presented in this chapter.

23.2 Construction Applications Plastic laminates, particularly fiber-reinforced composites, have been widely utilized in the aerospace applications and defense industries because of their light weight, high specific strength, high specific stiffness, and potential for very low maintenance cost over their life of operation [1–4]. Laminated composites, initially developed for the military aerospace market, offer performance comparable to that of conventional structural metals, and now find applications in communications satellites, aircraft, sporting goods, transportation, heavy industry and in the energy sector, in both wind turbine construction and oil and gas exploration [4]. The textile industry has developed the ability to produce net-shape/near-net-shape fabrics using highly automated techniques such as stitching, weaving, braiding, and knitting. Super-heavyweight nonwovens can significantly reduce the number of plies required for a layup, making fabrication more cost-effective, especially for large industrial structures [4]. Noncrimp fabric (NCF)-based composites have become an attractive alternative for aerospace, marine, and even automotive applications due to their excellent performance and relatively low cost [5]. Fabric prepregs are generally preferred on complex contours because of their better ability to drape and conform to the mold [6]. However, many industrial applications, including circuit boards, heat exchangers, electronics protection, etc., require an improvement in the thermal conductivity of plastics [7–10]. Thermally conductive materials are designed by blending polymeric matrices with thermally conductive fillers. Among typically-used fillers are graphite and metallic powders [11–14] as well as boron nitride [15] and synthetic diamond powder [16]. In order to obtain materials with the desired level of electrical conductivity, polymers are frequently blended with different kinds of fillers, because of the insulating character of most industrial plastics [17, 18]. Such materials (antistatic, semiconductive or conductive) are used in many practical applications as heating elements, temperature-dependent resistors and sensors, self-limiting electrical heaters and switching devices, antistatic materials for electromagnetic interference shielding of electronic devices, etc. [19–21]. This chapter deals with the range of reinforced plastic laminates used in advanced polymeric composite applications.

23.2.1

Automotive Applications

Fiber-reinforced polymeric composites have become an increasingly attractive alternative to metal for many automotive components. Composite material suppliers and manufacturers have spent many years developing and producing lightweight components for automobiles and other wheeled vehicles, aimed at improving fuel efficiency and cost [7, 8]. Composites were adapted earlier and have been used for a much longer time in

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many automotive components [9–18]. In the following sections, the polymer composites discussed are used or being considered for use in the automotive industries. The Carrera GT high-performance racing car developed by Porsche AG (Stuttgart, Germany), features carbon/epoxy composites (with some aramid) in the chassis, doors, doorsills, wheel wells, trunk lid, hood, underfloor, center console, and bucket seats (Figure 23.1). The chassis is a sandwich construction of woven carbon fiber prepreg with a honeycomb core made of Nomex aramid paper [19]. Most of the Porsche Carrera GT’s exterior panels are fabricated using autoclave and prepreg technology. Exceptions include the front side sections and the rear center section and closing panel, which are manufactured using resin transfer molding (RTM) because their shapes are less amenable to prepreg molding. It could produce weight savings of approximately 40 percent over conventional metallic construction. A high-performance composite automobile chassis of the Lamborghini’s Gallardo, Maserati’s Spider and Ferrari’s F430 is fabricated by Italian automotive composites specialist ATR Group (Colonnella, Italy) via resin transfer molding (RTM) (Figure 23.2). Carbon/epoxy prepreg was fabricated in matched metal molds in the RTM process. After the ATR chassis is assembled from RTM’d subcomponents, the entire structure is

Porsche Carrera GT Rolling Chassis

A-pillar structure with comolded steel tube roll bar

Engine frame (high-temp woven prepreg/aluminum honeycomb sandwich structure), weighs only 20.4 kg/45 Ib

Unidirectional reinforcement stiffens door sill Integrated carbon fiber console

Monocoque chassis (fabric prepreg/Nomex honeycomp sandwich structure) with integrated A and B pillars, weighs only 79.4 kg/175 Ib

Chassis/engine frame assembly bolts

Figure 23.1 Patent-pending carbon fiber monocoque chassis delivers maximum stiffness-to-weight ratio in high temperature environment.

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Figure 23.2 ATR’s ‘space frame’ chassis is an all-composite design that can be customized to a specific manufacturer’s car design. Co-cured metallic fastener attachment points are designed in to accommodate various car components.

autoclave cured to ensure that consolidation pressure and heat are sufficient to produce a monocoque structure [20]. Approximately 107 kg of carbon fiber epoxy is used on the ATR’s chassis, providing a weight saving of almost 35 percent compared to an aluminum design. The doors produced for the 2003 Dodge Viper SRT-10 incorporate advanced design and tooling techniques. Working with Meridian and Quantum Composites, DaimlerChrysler developed an insert join technique, overlapping the carbon fiber sheet molding compound (CFSMC) on each side of the low density glass SMC in the charge pattern [21]. The test results for the assembled door indicated a 206 percent improvement in door sag stiffness and a 350 percent improvement in permanent sag deflection. DaimlerChrysler has also developed the new Viper’s six-piece fender support system (6.14 kg, or 13.5 lb) to replace an assembly of 15 to 20 metal parts (24.2 kg, or 53.1 lb) [21]. The fender support constructed of random carbon fiber composites meets all design load and deflection criteria. Compression molds were used to cure carbon fiber–vinyl ester prepreg fender. Carbon fiber fender support system provides significant cost advantages in addition to weight savings. Low cost and damage tolerance make natural fibers attractive for automotive applications with low load requirements. According to FlexForm Technologies (Elkhart, Ind.), which supplies thermoplastic sheeting reinforced with plant fibers, natural fiber composites offer a lighter weight and more environmentally friendly alternative to wood flour-filled plastics, as well as 25 percent improvement in strength in applications such as door panels and inserts, package trays, head liners, seat backs, sidewalls, pillars and center consoles [22]. DaimlerChrysler introduced the application of natural fiber composites in a door panel, starting with the Mercedes M-Class SUV and R-Class Sport Tourer [23]. This panel was built of natural fiber composite material. The material composition is typically 50 wt% of polypropylene, and 50 wt% of natural fiber blend consisting of kenaf, hemp, and flax in various combinations. The door panel made with a natural fiberreinforced sheet is also used in a number of vehicles, including the Chrysler Sebring, Ford Freestyle and Expedition, Audi A8, Mitsubishi Space Star minivan, BMW 7-series, Jeep SUV (Figure 23.3), and Iveco heavy truck line. The rear cargo area load floor of the Porsche Cayenne was constructed of structural layers of natural fiber composites with an expanded polypropylene foam core and covered with a carpet cloth. All materials are comolded in a single, low-pressure press cycle [23]. Among the vehicles using these floors are the Volkswagen Touareg and new Audi Q7 sport utility vehicles. The spare wheel cover of the Mercedes A-Class is produced using the direct long fiber thermoplastic (D-LFT) compounding technology, in which

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weighs only 79.4 kg/175 lb

Figure 23.3 A door liner for this Jeep SUV is made with Flex Form natural fiber-reinforced thermoplastic sheeting, with fasteners bonded and integrated during processing.

raw polypropylene and abaca fibers are melt compounded in a twin-screw extruder, producing a molten charge that is subsequently compression molded in a cold tool [23]. The rear tub on the Pontiac Solstice was launched in 2006 and built using a preforming and press forming process developed by Molded Fiber Glass Co. (Ashtabula, Ohio) (Figure 23.4). The preforms contain 38 to 40 percent glass by weight in 2–3 inch fiber lengths [24]. When the preform is laid into a heated compression tool, a measured amount of polyester resin is distributed over the surface before closing the mold. An individual composite rear tub, with dimensions of 130 cm wide by 91 cm long and up to 61 cm deep, is 30 percent lighter than its conventional counterparts. An approximately 40 percent cost savings was achieved on the preformed rear tube compared to steel. Advantages of using composite also include an improved sound and vibration dampening. The 602 hp 2009 Chevrolet CorvetteZR1’s hood, fenders, roof, roof bow, lower rocker moldings and front splitter are all autoclaved carbon fiber composite. About 3.5 kg weight savings was achieved on the carbon roof panel and roof bow compared to the SMC versions used on other Corvette model 8s [25]. In 2005, a fleet of 30 fuel cell vehicles (FCVs) were deployed by city governments and research organizations for real world testing. A carbon fiber-reinforced plastic (CFRP) decklid produced for the Ford Focus FCV helps offset fuel-cell system weight to meet vehicle weight targets [26]. Components targeted for conversion include the outer panel, the inner panel, and local reinforcements for latch, hinges and rubber bumpers that mount on the inner panel (Figure 23.5). The assembled composite decklid, weighing 4.75 kg, replaced a multipiece steel decklid weighing 10.8 kg. The solid and sandwich CFRP constructions are processed via vacuum bag/autoclave on carbon composite tools using an in-mold primer (Figure 23.6).

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Figure 23.4 The rear tub on the Pontiac Solstice is produced in a pre-forming and press forming process developed by Molded Fiber Glass Companies. Pre-forms are manufactured in a robotic process known as ‘positive glass placement’.

Ford Focus FCV CRFP Decklid (interior view)

Vent opening Class A outer panel

1,260 mm/60 inches Carryover bumper stop

Vent port for hydrogen storage tank

Inner panel

450 mm/18 inches

430 mm/17 inches

Bonded and primed inner/outer assembly weighs 4.5 kg/9.5 Ib (60 percent mass reduction)

Solid Iaminate (thckness varies from 1.1 mm to 2.7 mm or 0.04 inch to 0.11 inch)

Aramid honeycomb cored sandwich construction (7.2 mm/0.28 inch thickness)

Latch reinforcement for carryover closure hardware

Figure 23.5 Find a materials/process solution that minimizes decklid weight and can be fabricated cost effectively in extremely low volumes yet accommodates carryover components and duplicates exterior appearance.

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871

The fully assembled composite decklid inner and outer, viewed from inside.

Composites showed evidence of continued penetration into critical engine components that incorporate a host of improved technologies, designed to ensure high reliability and minimal maintenance costs. Engine components such as the transmission torque converter, brake piston, commutator, pulley, and other components call for stiffness, strength, dimensional stability, chemical resistance, fatigue resistance and creep resistance under high loads at elevated temperature [27]. DaimlerChrysler saw an opportunity to replace magnesium valve covers on the 4.7L V-8 engine with glass reinforced bulk molding compound (BMC), thus improving sealing and noise performance and reducing cost [28]. The BMC valve covers, built by Dana Corporation (Paris, Tenn., U.S.A.), are molded in dual cavity molds via straight injection molding, which shortens the cycle time and makes the parts easier to deflash. In addition, while the magnesium molds have a life of 150,000 to 20,000 cycles, the covers show excellent performance with fatigue life of 1 million cycles when the composite materials are properly used in valve cover design. DaimlerChrysler has replaced the aluminum throttle bodies on several engines with injection-molded versions produced from a ‘zero-shrink’ BMC from Bulk Molding Compounds Inc. (W. Chicago, I11.) [25]. Composite throttle bodies, with both improved properties and increased suitability to automated processing, permit weight savings and cost competitiveness with traditional metal structures. The engine with the new throttle bodies is common to several platforms, including the Chrysler Pacifica, the Jeep Liberty and Grand Cherokee and the Dodge Dakota, Ram, Durango, Nitro and Viper. These are just a few of the composite applications that have been put in production. Other vehicle manufacturers have had similar experiences with varying degrees of success. All of these developments have shown the absolute need for simultaneous consideration of material, process, design, and manufacturing techniques.

23.2.2

Marine Applications

The commercial application of polymer matrix composites has generated significant interest in the marine industry, resulting in the initiation of many research and development programs. The literature on the design of reinforced plastics for use in marine environments is extensive [29–39], and this section includes materials that are particularly suited for the use of composites in loaded components. The Dutch company Conyplex BV (Medemblik, Netherlands) recently launched the Contest 50CS, a 50-ft hull yatch. Construction is PVC foam cored glass-reinforced styrene built on a skeleton mold made with tooling gel coat and polyester tooling resin (Figure 23.7). While initial efforts focused on the hull, the

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Figure 23.7 The construction procedure of Contest 50 CS yacht based on cored glass-reinforced styrene built on a skeleton mold made with tooling gel coat and polyester tooling resin.

one-piece deck is vacuum infused as well, using the same simple strategy [40]. The finished vacuum-infused 50-ft hull has been extracted from the mold and transported to another station for assembly with the deck. The infusion method resulted in a savings of 250 kg of resin on the hull alone compared to the previous sprayup process. Sea Ray (Knoxville, Tenn., USA) has built hulls for 18-ft to 24-ft sport boats and yachts, using the Robotic In-Mold Fiber Reinforcement (RIMFIRE) robotic preform technology [41]. Conformable glass fabric is laid into tight radii in the lower mold. After the mold is moved into the preform cell, the robot reaches into the mold, spraying chopped glass and heat-activated binder, and then deposits the glass/binder mix onto the mold surface. When the preform is completed, engineered fabrics are laid over the preform in areas of high stress, followed by overspray to trim off the flange of the lower mold. After the edges of the upper and lower molds are joined and sealed, resin is introduced into the mold via the zero injection pressure resin transfer molding (ZIP RTM) process, followed by a cure at 32 ◦ C for about one hour. The 56.4-m/185-ft yacht Scheherazade used carbon/epoxy composite to provide additional strength and rigidity in the masts, deck panel, booms and rigging components. The deck is made of carbon plies that make up approximately 12 percent (by volume) of its structure, which forms the top, with the hull providing the other ‘sides’ of an integrated structure functionally similar to a box beam [42]. Four layers of carbon/epoxy are laid on a layer of marine-grade plywood, followed by a Divinycell foam core, an additional three layers of carbon/epoxy and a teak decking. In order to achieve optimal strength at the designed weight, biaxial carbon fabrics wet out with epoxy, in a vacuum bagged, room temperature-cured layup. Another relatively unusual structural use of carbon is Scheherazade’s stemhead fitting assembly, which is integral to the forward part of the bow and attaches the stainless steel bushings head stay to the hull [42]. Four stainless steel bushings sit on an angled composite faceplate that is supported by two web faces and a neat epoxy fillet. The plate, web faces and bushing surround are built of unidirectional carbon/epoxy prepreg. Structures are optimized by choice of

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Figure 23.8

873

3DL composite sails fly on nearly all top-end yachts involved in racing and performance cruising.

fiber orientation to provide uniform safety. After vacuum bagging, composite structures most often employ 120 ◦ C curing systems and the bushings are bonded to the faceplate using a secondary wet layup of the straps which is made of epoxy-coated biaxial fabric. Modern production techniques and advanced design procedure in the laminated sailcloths have led to a steady and progressive development in performance [43–48]. The sail of the Genuine Risk 90 racing sloop was constructed of 3DL carbon yarn and scrim over a polyester base film which prevents wind from passing through the sail in use (Figure 23.8). After the base film is tensioned out over the mold, adhesive side up, fiber placement head applies resin coated carbon or aramid yarn onto the complex curvature pattern that follows the specified primary structural loads [49]. In the next station, the scrim which provides tear strength and handles secondary loads is tensioned out over the yarn layup, adhesive side down. Then, the top vacuum bag is built over the scrim and vacuum is pulled, which takes about 20 to 30 minutes. The sail is allowed to cure for about 10 to 20 minutes before it is pulled from the mold. Marine propellers made of hybrid carbon–glass fiber-reinforced plastic have a successful in-service record in the repetitive collapse of vapor bubbles on the blade surface, which creates unwanted noise and imparts localized stress, ultimately damaging the blades and reducing propulsive efficiency. The composite marine propeller with a 2.9-m/9.5-ft diameter, developed by QinetiQ (Rosyth, U.K.), consists of centermost plies of carbon fiber. The blade is stiffened with glass fiber-reinforced plastic [42]. It is claimed that the use of glass fabric-faced polyurethane coating isolates the carbon fiber from the hostile environment of seaspray with sand and provides some increased impact resistance on the blade surface. The composite blades show a 35 percent weight saving compared to thinner metal blades, reducing wear and tear on system components such as gears and bearings. The composite propeller was installed on QinetiQ’s 90-m/295-ft long RV Triton trimaran warship prototype and subjected to tests during sea trials. On the basis of gathered data of tests on blades, the composite design provides a smooth take-up of power and reduces vibration on the main shaft compared to its metal counterparts. The carbon-composite driveshaft was designed to warship standards and installed on the Expeditionary Fighting Vehicle (EFV), built by U.S. Marine Corps. The composite shafts are fabricated by McClean Anderson (Schofield, Wis., U.S.A.), in a continuous filament winding operation [42]. Such driveshafts most often employ at 175 ◦ C curing systems. This elevated-temperature cure provides most environmentallydurable composite, particularly in its strength and modulus retention after moisture exposure. A 70 percent

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Figure 23.9

Composites are an enabling technology as offshore platforms and vessels move to deep water.

weight saving has been achieved for the 1-m/3-ft long marine driveshaft with a diameter of approximately 15-cm/6-in manufactured from carbon fiber compared to steel driveshaft. Offshore productions of fiber-reinforced plastic have a successful in-service record in the hostile environment of seaspray with sand, which causes erosion and corrosion (Figure 23.9). Composites used in offshore facilities meet a design and certification requirements and offer significant weight advantages [50–53]. Because they readily adapt to innovative manufacturing techniques, composites also can provide significant cost reductions. Drilling risers for use at great water depths are subject to compression and possible failure because the longitudinal resonant period of the disconnected riser is close to that of typical wave periods. Spencer Composites Corp. (Sacramento, Calif., U.S.A.) have fabricated several full-scale, high-pressure drilling riser joints that have been successfully field tested on the Heidrun Platform in the North Sea [52]. About 40 layers of a high-strength carbon fiber tow are wet out with epoxy resin and wrapped around the titanium liner protected with a hydrogenated nitrile rubber coating. The carbon fiber is wound in hoops and low-angle helical layers. Next in the assembly line is the winding of E-glass fiber tows which are applied primarily to secure the outer surface of carbon reinforcement. The riser then heads to the joining facility where the joints are overwrapped with fiberglass outer layers wet out with epoxy. It has been reported that risers manufactured from carbon fibers have a burst pressure of 12,000 psi/827 bars and an axial load capacity of 1.361 million kg/3 million lb. The test results indicate that the bending stiffness of the composite joint – 195 MNm2 /6.79 × 1010 lb-in2 – had to be the equivalent of the titanium drilling joints normally used on the Heidrun Platform.

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It has been proposed that submarine pipelines could be constructed with circumferential carbon fibers for resistance to external pressure and longitudinal glass fibers for lengthwise flexibility [50]. Ameron International, Fiberglass Pipe Division (Houston, Texas, U.S.A.) has manufactured its Bondstrand Series 7000M and PSX-L3C piping for the ring main piping, together with PSX-JFC for the dry deluge piping downstream from the ring main [52]. This fire system was installed on the Spar platform in the Gulf of Mexico. Fabrication of the 7000M piping begins by filament winding E-glass roving on a mandrel to form a liner over which carbon fiber tow is wrapped. The liner and roving are wet out with an amine-cured epoxy resin. Next, the pipe is wrapped with Ameron’s PSX (polysiloxane-modified phenolic) resin matrix which gives off water and formaldehyde during cure and improves impact resistance of the finished piping. Finally, about 10 plies of polypropylene veil tape and glass roving are wrapped around the whole pipe. When the veil vaporizes under extreme jet fire conditions a thin sleeve of polypropylene tape leaves an air gap in the laminate that works in combination with the phenolic and the glass rovings to create a heat barrier. The case studies considered in the paragraphs above discuss just a few of the many ways composites are performing in offshore oil platform applications. Composite products are used in a number of components in offshore applications, including blast enclosures, heat shields, crane trusses, gratings, floors, ladders, railings, stairways, storage cabinets, valves, and wind walls. Carbon composites are finding new marine applications not previously possible, but volume remains small. Although marine applications are regarded as a strong market for glass-reinforced composites, after many years of development, the marine market appears ready for carbon composite production commitment. The future for carbon composites remains bright, and an ever-increasing application to new design is predicted.

23.2.3 23.2.3.1

Other Applications Composites in Sporting Goods

Specialized composites also are making significant impacts in the sporting goods industry [2, 3, 54–56]. Carbon fiber-reinforced composites are well-established in equipment for warm-weather sports, particularly in golf club shafts and tennis racquets, but they also have made inroads into winter sports, in both skins and in often-employed snowboards [54]. The rise of the composite golf shaft during the 1980s and 1990s powered a huge increase in the use of carbon fiber-reinforced polymer. A typical Penley (San Diego, Calif., U.S.A.) golf shaft is a thin-walled tube ranging from 114 cm/45 inches to 127 cm/50 inches in length with an average weight of only 50 g to 80 g (1.77 oz to 2.83 oz) [56]. Both the torsional and bending loads applied to golf club shafts, either in mishap or in hitting the ball, must be considered in golf club design. The manufacturing process begins with two helical wraps around the mold-release-treated mandrel, at an angle of +/–45◦ , to handle torsional and radial impact load, followed by several 0◦ axial wraps which increase the stiffness of the rod. The axial wraps are followed by hoop plies for further strength and controlled flex. Typically, carbon-boron/epoxy hybrid prepreg is used throughout the area just above the hosel, due to its small diameter and proximity to the club head where flex forces are high. When the layup is complete, a clear polypropylene shrink film is wrapped around the shaft (to consolidate the layup during oven cure), followed by a cure that starts at ambient and ramps up to 135 ◦ C, then back down. After curing for about two hours, the mandrels are hydraulically extracted. The use of composites to construct carbon fiber bike frame components was one of the extensive applications of modern composite technology in the sport markets [57]. In 2004, Trek Bicycles produced its lightest frame to date – a scant 2 lb – for the OCLV 55 Madone SSL, which was developed specifically for the 2004 Tour de France Stage 16 uphill time trial on the Alpe d’Huez – the most anticipated stage of the Tour [58]. Trek has developed an in-house bladder molding process called OCLV (Optimized Compaction, Low Void), to apply external and internal pressure and heat to carbon/epoxy prepreg under extreme conditions to yield a highly

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Figure 23.10 This fireman’s SCBA, designed by Hyper Comp and made by Draeger, features a 6061 aluminum liner fully overwrapped with pre-impregnated T700 carbon fiber tow and finished with a thin layer of E-glass for abrasion protection and clear gel coat for a smooth, easily cleaned surface.

consolidated hollow part with high fiber volume. A successful OCLV process depends on exact placement of each subsequent unidirectional prepreg to build a multiaxial fiber architecture for each part with staggered ply drop-offs to avoid stress risers. Once the layup is complete, preprogrammed CNC equipment cuts the stacks into mirror-image preform using standard steel dies. After the preforms are placed into two-part female molds made of P20 steel and prepped with release agent, a deflated bladder is placed between them. Next, the mold is closed around the bladder and inserted into a press and the bladder is inflated before cure begins. After cure at 121 ◦ C for about one hour, parts are demolded and transported to another station for machining. Attach points are reamed and holes are drilled for mechanical fasteners. Then, inserts are bonded in to strengthen parts and oven cured. In the next station, adhesively bonded tubes and lugs are assembled into a bike frame with the aid of a pneumatic press for accelerated cure at 65.5 ◦ C to 82 ◦ C. 23.2.3.2

Composites in Pressure Vessels

Winding with towpreg has the potential of making significant inroads into pressure vessel manufacture and may revolutionize the amount of composites and types of techniques used in commercial composite cylinder manufacture [59–61]. The fireman’s self-contained breathing apparatus (SCBA) (Figure 23.10), designed by HyPerComp and made by Draeger, was built using towpreg to wrap a seamless 6061 aluminum liner with pre-impregnated carbon fiber tow [62]. The fiber was wound in hoops layers at 90◦ to the horizontal, followed by subsequent hoop and helical layers. After all carbon layers are completed, an outer E-glass skin is wound for abrasion protection. The cylinder is then hung in an oven for vertical cure and finished with a thin layer of gel coat for a smooth, easily-cleaned surface. Burst examination of the wrapped composite cylinders indicated that they were able to demonstrate a pressure of 3.4 times their service pressure. Although some minor damage was induced in the cylinder during low-pressure and full service pressure cycling testing, the successful fabrication of these cylinders satisfied the burst strength requirement (3.06 times service pressure).

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Figure 23.11

23.2.3.3

877

Two long runs of filament wound Dualam pipe. Seamed joints are visible near the I-beam support.

Composites in Corrosion

Because the direct cost of repairing or replacing corrosion-damaged infrastructure, piping and industrial equipment is estimated to be in the hundreds of billions of dollars per year, in new applications, composites are increasingly finding a place in corrosion repair, particularly in the oil and gas industry in refinery pipelines that must withstand highly corrosive by-products of oil production [63, 64]. Dual-laminated pipe is shown in Figure 23.11. The multimandrel filament winding machinery and computerized pipe fusion equipment, designed by CPF Dualam Vancouver (Port Coquitlam, British Columbia, Canada), made it possible to fabricate and install 25,000 ft/7620 m of dual laminated-pipe at the new BASF chemical complex in Shanghai, China [65]. Pipes for this project ranged in diameter from 1 inch to 24 inches (25.4 mm to 609.6 mm) and features a thermoplastic liner tube fully laminated using corrosion-resistant thermoset vinyl ester resins, surface veils and glass chopped strand mat. An embedded woven glass fabric layer is required between the liner and the outer glass/vinyl ester layer to ensure a mechanical lock with the outer thermoset reinforcement, thus preventing concentration of stresses that could cause mechanical damage. Once the liner is embedded with the glass fabric, it is loaded into the multimandrel filament winder. Next, a last E-glass winding strand is filament wound with an epoxy vinyl ester for strength. The cylinder is then removed from the winder and a thin layer of exterior finish is used to provide an ultraviolet (UV) resistant and flame retardant. According to the company’s calculations, dual-laminate pipe provides a weight savings of approximately 75 percent and significant cost savings (estimated to be more than 50 percent) compared to the lined steel and titanium pipe, respectively.

23.2.3.4

Composites in Energy

In today’s highly competitive wind energy market, the use of composite turbine blades is extending because a lighter and/or more efficient blade decreases the demands on the hub components and tower structure, decreasing capital and operating expenses for the entire turbine [66–68]. Such was the case, for example, with the 36.8 m long glass/epoxy blades manufactured by LM Glasfiber (Lunderskov, Denmark), which were put into service on the offshore wind farm’s 2 MW turbines in Copenhagen, Denmark. According to the company’s designers, vacuum-assisted resin infusion molding (VARIM) was found to be more cost effective

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Figure 23.12

The aerodyn 56.5 m blade nears completion.

and to provide better control over the materials than prepregs, especially for larger blades [67]. In order to damp the edgewise vibrations of the blade and minimize fatigue loads, a wide band of structural fiberglass composite was incorporated into the layup, which extends from the tip of the blade along the leading edge about two-thirds of the length of the blade. Composite blade manufacturers take a long look at carbon/glass hybrids, as building lightweight blades of greater length remains a primary focus for utility-size turbine manufacturers [68]. While an optimized glass-only design was sufficient for its 54m/176 ft long 54P offshore model, LM has introduced carbon fiber in the highly loaded areas of its 61.5 m/200 ft hybrid blade for a 5 MW turbine. The 56.5 m/185 ft carbon/glass hybrid blade with a weight of 15,700 kg/34,540 lb (Figure 23.12) was built by aerodyn Energiesysteme GmbH (Rensburg, Germany) for a 5 MW turbine, which will be deployed in the North Sea as part of Germany’s first deepwater offshore project. The use of carbon fiber in the 44 m V90-series blade, fabricated by Vestas Wind Systems A/S (Ringkøbing, Denmark), was claimed to reduce bending of the blades because of carbon’s greater stiffness. A preliminary cost benefit study indicated that significant weight savings [68] could be achieved by using carbon/glass hybrids in the blade constructions mentioned above. 23.2.3.5

Composites in Infrastructure

Recently, researchers have focused considerable research and development efforts on fiber-reinforced plastic bridge decks due to growing recognition of composite’s durability, corrosion resistance and light weight, and how those properties translate to fast installation, better load-carrying capability and virtually no maintenance over the life of the bridge [69–72]. The 22-ft long and 32-ft wide three span bridge (Figure 23.13), featuring a new pultruded fiber glass/vinyl ester deck, was fabricated by Martin Marietta Composites (Raleigh, N.C., U.S.A.) and installed in Greene Country, Ohio, as part of the National Composites Center’s (NCC) Composites for Infrastructure (C4I) initiative [72]. The actual live load testing with an array of strain gauges was conducted using a fully loaded truck weighing 34.5 tons. The test results indicated that most of the data were less than 40 percent of the allowable 0.118-inch deflection under current industry guidelines for the bridge’s size and traffic load. Four different composite deck types, from four different manufacturers, were installed on the 207 m/ 673 ft-long Salem Ave. Bridge in Dayton, Ohio [72]. The deck, consisting of pultruded trapezoidal profiles,

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Figure 23.13 This composite bridge deck was installed in Greene Country, Ohio, as part of the National Composites Center’s (NCC) Composites FOR Infrastructure (C4I) initiative.

was manufactured by Creative Pultrusions Inc. (Alum Bank, Pa., U.S.A.) and then adhesively bonded together to achieve the specified deck panel lengths for a job. The hardcore segment was a resin infused, cored structure with fiberglass fabric-wrapped foams and fiberglass skins, fabricated by Hardcore Composites. The segment, devised by Infrastructure Composites International (ICI, San Diego, Calif., U.S.A.), features a core which was made with chopped strand mat with sinusoidally shaped webs to create very strong core materials. The fourth segment was built by Composite Deck Solutions LLC (Dayton, Ohio, U.S.A.) as a hybrid design made up of pultruded fiberglass stay-in-place forms filled with high-strength concrete, with internal fiberglass reinforcing grids. For the Salem Ave. Bridge, the panels were approximately 2.4 m/8 ft wide, 12.2 m/ 40 ft long and 216 mm/8.6 inches thick, and all were coated with a thin weather-resistant polymer concrete wearing surface. Adjacent panels were attached to the underlying steel girders with a combination of epoxy adhesive and studs welded to the top flanges of the girders. After the panels were tied together, holes drilled in the panels to accept the studs were filled with high strength grout. The results of live load tests performed with the aid of fully loaded trucks indicate that the maximum deflection at mid-span was considerably less than the American Association of State Highway and Transportation Officials’ (AASHTO) limitation of 8.8 mm/0.35 inch.

23.2.3.6

Composites in Ballistic Protection

Trends clearly indicate that the use of fiber-reinforced composites will continue to grow in antiballistic systems, supplanting and supplementing legacy systems that rely on metals and ceramics [73–76]. Honeywell Advanced Fibers and Composites Inc. (VA, U.S.A.) offers a bullet resistant vest, Spectra Shield, which is a flexible, cross-plied (0◦ /90◦ ) nonwoven fabric that enables more fibers to engage the round, dispersing a bullet’s energy 360◦ throughout the panel for outstanding stopping power [75]. Nonwoven fabric, consisting of Honeywell’s patented polyethylene Spectra fiber, is impregnated with thermoplastic resin to optimize its impact absorption characteristic (Figure 23.14). According to Honeywell, Spectra fibers exhibit a tensile modulus between 900 g/d (grams per denier) and 1500 g/d, compared to a tensile modulus of less than 500 g/d for steel, E-glass and S-glass. Spectra Shield is combined with plates consisting of a ceramic strike face with a Spectra Shield spall liner. The ceramic helps break up a bullet or turn it on its side to diffuse its impact force and rob it of its velocity, while the spall liner produced by Honeywell’s polyethylene Spectra fiber deals with over-matching threats and reduces damage caused by fragments.

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Figure 23.14 Spectra Shield from Honeywell is a flexible, cross plied (0/90) nonwoven fabric used in soft armor vests. It is impregnated with thermoplastic resin to improve its ability to absorb impact.

Another notable application of a composite material is a blast-resistant laminate called ThermoBallistic line, developed by Polystrand (Montrose, Colo.), made from cross-plied (0◦ /90◦ ) nonwoven fabric reinforcement in a matrix of Polystrand’s proprietary polypropylene [75]. Since the establishment of NIJ Standard 0108.01 Level IIIA Requirement of V0 (1400 ft/sec, ± 50 ft/sec) threat level, which is equivalent to the threat posed by a 9 mm handgun, a 3.4 lb/ft2 ThermoBallistic-E made with E-glass fiber is required, while ThermoBallistic-S, the company’s S-glass version, offers the same protection at 2.5 lb/ft2 . The company’s products generally consist of a ceramic strike face with a composite backing when armor must adequately diffuse a greater threat. The composite backing is made of Kevlar, produced by DuPont (Richmond, V.A.), which is one of the most commonly-used fibers in vehicle armor systems facing improvised explosive devices (IEDs) threats because it is inherently nonflammable, offering high temperature resistance at continuous temperatures as high as 150 ◦ C. A notable example of armored vehicles used Polystrand’s material is the M915 Tractor Cab from the ArmorWorks (Tempe, Ariz.) for the U.S. military’s over-the-road truck. The M915’s armor provides substantial protection against high-velocity small arms fire, improvised explosive devices and roadside bombs.

23.2.3.7

Composites in Mass Transit

High-speed train technologies have become practical and cost-competitive, in large part through the use of composites to noses and other components. In addition to meeting mechanical property and processibility

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Figure 23.15 Composite noses are attached to Acela locomotives. These contoured components not only provide smooth aerodynamic performance but must meet tight geometric tolerances as well.

requirements, all materials used within the pressurized portion of the high-speed train must meet flammability resistance requirements defined by regulatory agencies and, if applicable, smoke and toxic-gas emission guidelines of the train manufacturers. An example of such a pressurized portion using polyester composites is the nose of the Amtrak’s Acela Express (Figure 23.15), which carries passengers at up to 240 kph, travelling from New York City to Washington, D.C. in less than three hours [77]. The predominant fiber used is fiberglass; however, balsa wood core was employed throughout the unsupported body of the nose to attain the required stiffness. The Acela’s noses – measuring approximately 3.6 m high by 3.2 m wide by 4.6 m long – not only provide smooth aerodynamic performance but also meet tight tolerances of the rest of the train, which requires a perfect fit. Incorporating alumina trihydrate (ATH) into a polyester resin has been found to be an effective way of improving flame retardance of the Acela’s exterior nosecone. Other aerodynamically contoured composite structures are the Acela’s shrouds that surround the two roofmounted air conditioning units per coach car to control both noise and air inlet volume. The Acela’s shrouds were made using hand layup and vacuum bagged with fiberglass/polyester resin. The equipment mounted on the underside of the train car are also protected by composite cover. Siemens’ Velaro E and ICE trains use composite undercarriage covers because of increased aerodynamic demands as well as greater impact from rocks kicked up by the train. The covers that generally employ fiberglass/polyester around a balsa core were fabricated using hand layup and vacuum bag molding techniques. The polyester resin used for aerodynamic shrouds and undercarriage covers is required to have some degree of resistance to burning, which can be accomplished by adding alumina trihydrate to resins, depending on the resistance required. Composites are also widely used in the interiors of high-speed trains. The predominant design considerations for interior components are impact resistance, stiffness, surface smoothness and warmth and noise insulation to satisfy the train journeys and their customers. Interior components such as luggage racks, sidewalls, ceilings, galleys, lavatories and cabinet coverings are routinely made of glass fiber-reinforced phenolic sandwich construction with a honeycomb core. Seat bowls are formed in heated compression molds from fiberglass/polyester SMC. To ensure excellent fire-resistant properties, including low flammability and low smoke and toxic gas emissions, the neat resin is filled with ATH prior to processing [78]. This section features fruitful areas of applications of laminated composite materials that have been more widely accepted in the automotive, marine and construction industries [79–84]. The notion of producing the lightweight components and reducing the manufacturing costs embraces the development of modern

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manufacturing techniques for composite production with improved quality and reliability. The future for polymeric composites remains bright, and an ever-increasing application to new design is predicted.

23.3 Aeronautics and Spacecraft Applications 23.3.1

Aeronautics Applications

Among the first uses of modern composite materials was about 30 years ago when boron-reinforced epoxy composite was used for the skins of the empennages of the U.S. F14 and F15 fighters. Initially, composite materials were used only in secondary structures, but as knowledge and development of the materials has improved, their use in primary structures such as wings and fuselages has increased. Initially, the percentage by structural weight of composites used in manufacturing was very small, at around 2 percent in the F15, for example. However, the percentage has grown considerably, through 19 percent in the F18 up to 24 percent in the F22 [85]. The AV-8B Harrier GR7 has composite wing sections and the GR7A features a composite rear fuselage. Composite materials are used extensively in the Eurofighter: the wing skins, forward fuselage, flaperons and rudder all make use of composites. Toughened epoxy skins constitute about 75 percent of the exterior area. In total, about 40 percent of the structural weight of the Eurofighter is carbon-fiber reinforced composite material. Other European fighters typically feature between about 20 and 25 percent composites by weight: 26 percent for Dassault’s Rafael and 20 to 25 percent for the Saab Gripen and the EADS Mako.The B2 stealth bomber is an interesting case. The requirement for stealth means that radar-absorbing material must be added to the exterior of the aircraft with a concomitant weight penalty. Composite materials are therefore used in the primary structure to offset this penalty. The use of composite materials in commercial transport aircraft is attractive because reduced airframe weight enables better fuel economy and therefore lower operating costs. The first significant use of composite material in a commercial aircraft was by Airbus in 1983 in the rudder of the A300 and A310, and then in 1985 in the vertical tail fin. In the latter case, the 2000 parts (excluding fasteners) of the metal fin were reduced to fewer than 100 for the composite fin, lowering its weight and production costs. Later, a honeycomb core with CFRP faceplates was used for the elevator of the A310. Following these successes, composite materials were used for the entire tail structure of the A320, which also featured composite fuselage belly skins, fin/fuselage fairings, fixed leading- and trailing-edge bottom access panels and deflectors, trailing-edge flaps and flap-track fairings, spoilers, ailerons, wheel doors, main gear leg fairing doors, and nacelles. In addition, the floor panels were made of GFRP [86]. In total, composites constitute 28 percent of the weight of the A320 airframe. The A340-500 and 600 feature additional composite structures, including the rear pressure bulkhead, the keel beam, and some of the fixed leading edge of the wing. The last is particularly significant, as it constitutes the first large-scale use of a thermoplastic matrix composite component on a commercial transport aircraft. Composites enabled a 20 percent saving in weight along with a lower production time and improved damage tolerance. The A380 is about 20–22 percent composites by weight (Figure 23.16), and also makes extensive use of GLARE (glass-fiber-reinforced aluminum alloy), which features in the front fairing, upper fuselage shells, crown and side panels, and the upper sections of the forward and aft upper fuselage. GLARE laminates are made up of four or more 0.38 mm (0.015 in) thick sheets of aluminum alloy and glass fiber resin bond film. GLARE offers weight savings of between 15 and 30 percent over aluminum alloy along with very good fatigue resistance. The top and bottom skin panels of the A380 and the front, centre and rear spars contain CFRP, which is also used for the rear pressure bulkhead, the upper deck floor beams, and for the ailerons, spoilers and outer flaps. The belly fairing consists of about 100 composite honeycomb panels. The Boeing 777, whose maiden flight was 10 years ago, is around 20 percent composites by weight (Figure 23.17),

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Figure 23.16

883

Airbus A 380 ‘Super Jumbo’.

with composite materials being used for the wing’s fixed leading edge, the trailing-edge panels, the flaps and flaperons, the spoilers, and the outboard aileron [86]. They are also used for the floor beams, the wing-to-body fairing, and the landing-gear doors. Using composite materials for the empennage saves approximately 1500 lb in weight. The Boeing 7E7 will leverage extensive use of composite materials (estimates are as high as 50 percent) in the quest for very high efficiency and performance with reduced weight. The excellent strength-to-weight ratio of composites is also used in helicopters to maximize payloads and performance in general. Boeing Vertol used composites for rotorcraft fairings in the 1950s and made the first composite rotor blades in the 1970s. Composites are used in major structural elements of many modern helicopters, including the V22 tilt-rotor aircraft, which are approximately 50 percent composites by weight. The formability of composites has been used to particular advantage in helicopter manufacture to reduce the number of component parts and therefore cost. There are many parameters which must be examined

Figure 23.17

Boeing 777–200.

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Figure 23.18

Lockheed Martin F-22A Raptor JSOH.

when choosing materials for any specific project; these involve material and manufacturing costs, aircraft performance, material performance and costs to the user such as damage tolerance, maintenance costs, etc. The way in which these parameters are interwoven will depend considerably on the type of project being undertaken. A light aircraft for the Tiger Club would require a different mix of criteria to an advanced combat aircraft for the RAF. These in turn would be different from the selection that would be made for a one-shot vehicle such as a booster rocket. In order to narrow this field we will concentrate on the type of aircraft typical of Kingston designs, the small high-performance aircraft for use by air forces around the world. Usage of composite materials in aerospace applications has gradually increased as manufacturing technology matures and confidence in the material’s long term behavior improves [87]. For example, more than 25 percent of the structure of the F-22 Raptor is composites (Figure 23.18), while next-generation commercial aircraft such as Boeing’s are being designed with more than 50 percent of the aircraft’s structural weight, and all of its exposed surfaces, fabricated from composite materials. Today, the aerospace industry views carbon and glass fiber-reinforced composites as mature, viable competitors for metal alternatives [88]. The initial and still important motivation for aerospace industry application of composite material in place of metal is performance enhancement via weight savings, often in combination with other advantageous material-related behaviors such as low thermal expansion and/or reduced radar cross section. Interestingly, while it can often be shown that there is significant life-cycle cost advantage associated with the selection of a composite structure in place of an equivalent metal design, this factor is often of relatively low importance in material selection decisions since the acquisition and operational budgets usually come from different and unrelated pools of money. The growing usage of composites by both the military and commercial aerospace industry indicates that, at least for these traditionally weight-sensitive applications, composite material and process engineering has done a good job of keeping production costs of high-performance composite structures competitively low. The weight savings and mechanical strength and stiffness properties of composites are equal or greater importance in aerospace structure uses. These advantages led to early use of composite in space structures. Some of the general areas in which composites have been used in space include trusses, platforms, pressure vessels, tanks and shells. An additional requirement for materials in space applications is that they must withstand the hostile environment of space. This environment includes extremes of temperature, radiation and concentration of molecular oxygen.

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CFRP (carbon fiber-reinforced plastics) fuselage structures are expected to be realized with future generations of aircraft [89]. Going into the post-buckling regime with these structures requires improved fast and reliable procedures for analysis and design of stiffened fiber composite panels. The EC project POSICOSS, which was started in the year 2000, develops such procedures. A European consortium is evaluating materials for the construction of a new supersonic transport aircraft that may replace Concorde. Current designs propose to use an aluminum alloy for the fuselage which is required to have superior creep resistance and damage tolerance to the Concorde alloy 2618A [90]. Promising results have been obtained with extruded Al-Cu alloys containing minor additions of magnesium and silver which stimulate hardening by the relatively stable  precipitate. Data are presented showing that these alloys have tensile and accelerated creep properties which are better than those of competing commercial alloys of the 2000 series, together with satisfactory levels of fracture toughness. Of the four experimental alloys studied, the optimal composition is Al-5.6Cu-0.45Mg-0.45Ag-0.30Mn-0.18Zr (wt%). Carbon fibers/polymer matrix composites tend to be used more widely instead of aluminum structures in the aircraft and aerospace industry [91]. There are many reasons that explain the increasing interest in this class of composites due to their light weight, high strength, high stiffness, good fatigue life, excellent corrosion resistance and low cost manufacturing. Moreover, a considerable effort is paid in improving the thermal/electrical conductivity and the electromagnetic shielding effectiveness of lightweight composites. It is well known that, in operation, the engines of the aircraft generate a large amount of electrical current during start up. This electrical current is undesirable and must be conducted away. As well, the electronic components need to be protected against electromagnetic waves produced by other sources placed inside or outside the aircraft. The defects and damages in composite aircraft structures, especially honeycomb sandwich constructions, have been studied by means of an instrumented tap test and imaging system. The resulting tap test images based on the impact duration and displayed in a C-scan format can readily reveal the shape and extent of honeycomb core damages [92]. The images also reveal considerable detail of internal substructures such as core splices, ply build-up, and changes of foam core density. Based on a grounded-spring mechanical model, the tap scan images can be converted into images that show the quantitative changes of the local stiffness. A manual tap scan imaging system was assembled for laboratory experiments. Using actual composites parts from airplanes, impact damages and substructures were imaged. In addition, effects of the tapper mass, impactor radius, tap velocity and operator dependence were studied. A compact fieldable system is being built and an automated tapping and imaging system is also under development. Safety in aeronautics could be improved if continuous checks were guaranteed during the in-service inspection of aircraft [93]. However, until now, the maintenance costs of so doing have proved prohibitive. For this reason there is a great interest in the development of low cost nondestructive inspection techniques that can be applied during normal routine tests. The analysis of the internal defects (not detectable by a visual inspection) of the aircraft composite materials is a difficult task unless invasive techniques are applied. The analysis of the time/space variations in a sequence of thermographic images allows the identification of internal defects in composite materials that otherwise could not be detected. A neural network is trained to extract the information that characterizes a range of internal defects in different types of composite materials. After the training phase the same neural network is applied to all the points of a sequence of thermographic images. The experimental results demonstrate the ability of the method to recognize regions containing defects but also to identify the contour regions that cannot be associated either with a defective or with a sound region. Microsensors and microelectromechanical systems (MEMS) are currently being applied to the structural health monitoring of critical aircraft components [94]. The approach integrates acoustic emission, strain gauges, MEMS accelerometers and vibration monitoring devices with signal processing electronics to provide real-time indicators of incipient failure of aircraft components with a known history of catastrophic failure

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due to fracture. Recently, a combination of the need for safety in the air and the desire to control costs is encouraging the use of in-flight monitoring of aircraft components and systems using lightweight, wireless and cost-effective microsensors and MEMS. An in situ aircraft structural health monitoring (ASHM) system, with sensors embedded in the composite structure or surface-mounted on the structure, would permit the timely detection of damage in aircraft. In this section we give an overview of microsensors and MEMS (of physical dimensions of the order of a centimeter or less) and their associated driving electronics for health and condition monitoring of existing (and aging) and future aircraft and composites. Silicon micromachining offers the potential for fabricating a range of microsensors and MEMS for structural applications including load, vibration and acoustics characterization and monitoring. Such microsensors are extremely small; they can be embedded into structural materials, can be mass produced and are therefore potentially cheap. Additionally, a range of sensor types can be integrated onto a single chip with built-in electronics and ASIC (Application Specific Integrated Circuit), providing a low-power microsystem. Smart sensors are being developed using standard microelectronics and micromachining in conjunction with novel Penn State smart electronics or wireless communication systems suitable for condition monitoring of aircraft structures in-flight. The main application areas of this investigation include continuous monitoring of (a) structural integrity of aging aircraft, (b) fatigue cracking, (c) corrosion, (d) deflection and strain of aircraft structures, wings, and rotorblades, (e) impact damage, (f) delamination, (g) location and propagation of cracks, and (h) the quality of conventional bonds and ‘kissing bonds’ in composite structures. During the last two decades advanced thermoplastic composites with excellent material properties and good manufacturability characteristics have been developed. Since the early 1980s, widespread use of these materials in the aircraft industry has been anticipated. However, the expected breakthrough has not been realized; despite the intensive efforts that have been undertaken, application of thermoplastic composites in aircraft structures has been rare [95]. The key factor when introducing new materials and processes in series production is their cost efficiency; for producing aircraft components, cost efficiency has been proved to be lower when using thermoplastic composites than when using other more conventional materials such as aluminum products or thermosetting composites. At present, from this statement might be excluded only components that can be manufactured using press improvements in cost efficiency can be made when reducing the component manufacturing costs; depending on the case, manufacturing costs may contribute up to 80% of the total cost of the component. In this work, results of an on-going investigation for producing thermoplastic composite components cost-efficiently are presented. A concept for optimizing thermoplastic composite manufacturing processes with regard to component quality and cost is introduced; it is applied on the paradigm of the diaphragm forming technique. Application of this concept has driven the development of a new cold diaphragm forming technique for producing thermoplastic composite cost efficiently; the new technique will be also introduced. Thermoset adhesive, FM 73, is widely used in the bonded repair of metallic aircraft structures. The results at room temperature and at an elevated temperature show that the time-dependent properties of this film adhesive are significant even at room temperature. Considerable stress relaxation was observed during strain hold experiments and the adhesive was also found to demonstrate significant creep characteristics during creep tests [96]. The monotonic shear stress/strain behavior and time-dependent characteristics can be adequately described using a unified theory of plasticity. Results from a finite element analysis of an adhesively bonded joint show that significant error in the analysis may result during monotonic loading if the rate dependency of the adhesive is not taken into account. This has significant implications in the design and analysis of bonded joints. The advantage of low density of magnesium alloys in fiber metal laminates (FMLs) as compared to current light alloys comes along with various disadvantages in mechanical properties. To determine whether this alloy is applicable in an FML configuration for aircraft structures, the effect of the properties on the overall FML behavior has been addressed. An initial evaluation of FMLs based on current magnesium alloys,

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using knowledge and prediction models developed for current FML has been validated with a limited amount of data obtained from the literature. Emphasis in United States research on electric launch science and technology continues to focus primarily on understanding the fundamental physics and enabling technical issues for achieving hypervelocities. Programs to launch large structures (aircraft and space vehicles) at moderately low velocities also continue, mainly addressing the engineering challenges [97]. Although research topics in electric launch remain much the same as in past years, the level of detailed modeling and experiments of complex electromagnetic systems has increased dramatically. The U.S. Army is aggressively pursuing the engineering of an advanced electrothermal chemical (ETC) launcher but, in parallel, is making a strong computational and experimental effort to determine the fundamental interactions occurring at the plasma/propellant interface. For electromagnetic launchers, most of the effort is directed toward improved computational tools, exploitation of these tools for detailed understanding of transient electrodynamic phenomena, novel diagnostics, and experiments to resolve remaining critical issues such as transition from solid to arc contacts in railguns, improved computational techniques for pulsed power systems, and application of these tools to design new high-energy pulsed-power sources. New methods of testing and determining the critical properties of advanced materials, such as composites, are being developed to enable these materials to be evaluated in extreme thermal and electromechanical environments. Additionally, the U.S. Navy is also in the process of initiating hypervelocity electromagnetic launch efforts for extremely long-range artillery systems employing high-G novel projectiles. Other applications of electric launch technology, such as hypervelocity powder deposition and electromagnetic gun launch to space, continue to offer new and interesting opportunities. Design/methodology/approach presents a description of the automated tape layer (ATL) process and the fiber placement (FP) process [98]. These processes are the most ‘automated’ of all processes being used to fabricate composite aircraft structure. FP machines and ATLs are composites machine tools and are the closest comparison the composites industry has to metals machining equipment. There is a need for more variety of composites automation and more affordable machines in the aerospace composites industry. ATL and FP are composites laminating technologies that could be adapted to a wide range of machine sizes, configurations, and price ranges. More widespread use of automated processes in composites would tend to lower the cost of composite aircraft structure on a global basis. The application of composites in the aircraft industry has increased significantly over the past few decades. With traditional composite laminate shaping, each layer is made to conform to the mold surface by hand before subsequent layers are added. This is a very labor- and time-intensive process. There is a great deal of interest in developing an automated process for forming composite parts with compound curvatures [99]. The proposed composite forming process utilizes a computer-controlled, reconfigurable discrete element mold to incrementally form a compound curvature part shape from a flat layup, thereby facilitating process automation. An elastomeric interpolating layer is placed on top of the hemispherical forming ends of the die elements to prevent dimpling of the composite layup. The process employs vacuum to pull a single diaphragm (top), composite, and interpolator into contact with the mold surface. Through an experimental investigation, this new composite forming process with active tooling has been successfully demonstrated. Heating of the composite is accomplished by uncontained, forced convection using a matrix of heated air jets mounted above the composite. However, low-powered conduction is shown to be the best heating method in terms of both composite heating time and minimization of through-thickness temperature. Using vacuum to conform both the composite and the interpolator to the mold, and choosing sufficiently stiff diaphragm and interpolator materials, undimpled and wrinkle-free composite parts have been formed in an incremental fashion. For the production of blister fairings that form part of the thrust-reversing unit of a Rolls Royce RB 211 propulsion system, a double-diaphragm forming technique was used suitable for processing continuous fiberreinforced thermoplastic prepregs [100]. Within the optimization of the forming process, differences between draping and deep-drawing were investigated with respect to achievable finish quality of the components.

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Utilizing a grid strain analysis and a surface-fitting program, the blank shape and fiber architecture of the laminates were optimized prior to forming. Employing Plytron and Tepex prepregs a series of prototypes was manufactured and tested, first on a laboratory scale and later on a civil aircraft. For assessment of the feasibility of both the diaphragm-forming technique and the manufactured components, activity-based costing was applied to compare the glass fiber-reinforced thermoplastic parts with their currently used carbon fiber-reinforced thermosetting counterparts. An affordable approach to substitute metallic honeycomb panels with more durable stiffened graphite/epoxy composite panels has been demonstrated on a honeycomb panel for the Royal Australian Air Force (RAAF) F-111 aircraft [101]. This panel was selected with the aim of validating the design, manufacture and certification approaches. Future work will focus on airworthiness certification, operational testing, addressing service-related issues, and composite substitution of other types of metallic components. Dassault Aviation has been using composite materials for its business jets since 1979, when the Falcon 50 was fitted with composite ailerons. A more significant milestone came in 1985, when the first all-composite wing for a civil aircraft, and the first to achieve FAR Part 25 certification, was joined to a Falcon 10 fuselage, remaining in service for 20 years [102]. The company has also made full use of its many years of combat aircraft experience. Its Rafale omni-role fighter, for example, contains 30% carbon composites, while the new Falcon 7X, with its all-CFRP vertical fin and horizontal tailplane, has around 20%. By comparison, the new Airbus A350 XWB and Boeing 787 will contain 50% by weight of CFRP. Today, Dassault Aviation is working with European partners on two major development efforts: The Advanced Low Cost Aircraft Structures (ALCAS) programs, led by Airbus UK and Dassault Aviation, include development of an all-composite business jet fuselage and wing; and the Anglo-French-German full-barrel composite fuselage (FUBACOMP) programs, in which Dassault Aviation and BAE Systems are manufacturing a scaled version of a Falcon business jet front fuselage. This was one of the first major validations of fiber placement technology in Europe. Composites become more attractive for business jets, when they can be produced as a single, integrated ‘one-shot’ structure. The Falcon 7X broke new ground at Dassault Aviation with its one-piece aluminium upper wing panel, in which the entire top surface, complete with stringers, was cut from a single billet of metal and then creep-formed. The cost picture becomes better when it can be manufactured with composites. A durability and damage tolerance criterion drives the design of most composite structures. Those criteria could be altered by developing structure that repairs itself from impact damage [102, 103]. This is a technology for increasing damage tolerance for impact damage. Repaired damage would enable continued function and prevent further degradation to catastrophic failure in the case of an aircraft application. Further, repaired damage would enable applications to be utilized without reduction in performance due to impacts. Selfrepairing structures are designed to incorporate hollow fibers, which release a repairing agent when the structure is impacted so that the repairing agent will fill delaminations, voids and cracks in less than one minute, thus healing matrix voids. The intent is to modify the durability and damage tolerance criteria by incorporation of self-healing technologies to reduce overall weight: The structure will actually remain lighter than current conventional design procedures allow. The damage can be repaired to within 80–90% of original flexural strength in less than one minute, in laminates that are processed at 300–350F typical for aircraft composites. The main focus at Natural Process Design, Inc. was on testing of elements in compression after impact and a larger component in shear. The results show potential; with self-repairing composites, compressive strength is maintained sufficiently so that less material can be used as per durability and damage tolerance, yielding a lighter structure. Airframe structures consist essentially of an assembly of simple elements connected to form a load transmission path. The elements, which include skins, stiffeners, frames and spars, form the major components such as wings, fuselage and empennage. These connections or joints determine the structural efficiency of the airframe as they are potentially the weakest points. The tests [105] are concerned mainly with joints used to connect structural elements made of advanced fiber composite laminates – mainly graphite/epoxy

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(gr/ep) – to other composite parts or to metals, with bonded and mechanical joints used in the manufacture of airframe components and with joints required to repair structural damage. In many ways, the topic of repairs is the more challenging because in a repair situation the design and the materials options are usually very limited. Finite element analysis procedures are used in most practical designs. The materials aspects covered will be those essential to the manufacture of sound joints and repairs. Various hygroscopic effects of such parameters as hygrothermal temperature, matrix volume ratio (Vm ), void volume ratio (Vv ), specimen thickness, layup sequence and internal stress were investigated for epoxy/carbon fiber composite laminates by Choi and co-workers [106]. The specimen thickness and layup sequence had little effect on the through-the-thickness water absorption behavior of composite laminates, but the other parameters affected the moisture absorption rate and equilibrium water uptake in different ways and intensities. The glass transition temperature of composite laminates was strongly affected and linearly decreased by the quantity of equilibrium water uptake. A characteristic length of moisture migration through the unidirectional laminates was proposed as a function of fiber angle to the exposed laminate surface. In this approach, the fibers imbedded in the matrix were assumed to act as a burner to the penetrating water molecules, and the developed model compared well with the experimental results. Nowadays, aircraft manufacturers are not only looking for ways to reduce the structural weight of their aircraft but they are also searching for structural concepts that will lead to a cost reduction [107]. One way to realize a cost reduction is to design a component with a high level of part integration since this will lead to a reduction in labor-intensive trimming and assembly costs. By using composites in combination with new fabrication concepts this part integration becomes feasible. One of these new fabrication concepts is resin transfer moulding (RTM) with prepregs. In traditional RTM processes, dry fiber preforms are positioned in a mold cavity. After the mold is closed, resin is injected into it and the fibers are impregnated. In the RTM process, parts of the dry preform are replaced by prepreg. After closure of the mold, the mold is heated and the resin in the prepreg starts to melt. Then RTM resin is injected into the mold. The pressure of the RTM resin is used to pressurize the prepreg. The main advantage of this fabrication concept is that sub-preforms can be made very easily in prepreg that would be very difficult to make with dry fabric due to the lack of tack. Another advantage is that with the RTM process time is reduced, because only a small quantity of resin has to be injected. In order to demonstrate the feasibility of this fabrication concept, a hat-stiffened cargo door concept was developed. Two doors were made and were tested by applying a pressure difference to the door of 0.12 MPa. Neither door failed during the tests. Thanks to their low density, good thermal, mechanical and tribological properties, composites made of carbon fibers and carbon matrix (C/C) are particularly adapted to the manufacture of aircraft brake discs. Several methods have been developed to improve their performance [108]. For instance, the introduction of silicon carbide within the matrix notably increases the friction coefficient. While carbon fibers enhance the toughness of the matrix, silicon carbide shows high hardness, thermal stability and low chemical reactivity, leading to superior friction properties. The purpose of the present study [108] was to evaluate the influence of such matrix modification on the friction behavior and to identify the related mechanisms. With C/C as a reference, four different hybrid matrix composites were elaborated. For the first one, the matrix consisted of carbon (C) and silicon carbide (SiC), deposited in alternate layers around the fibers. The second was constituted of C, SiC and free silicon (Si), randomly distributed. For the third one, the matrix was made of SiC onlyThe fourth was similar to the third but with some additional free Si. These samples were submitted to structural and mechanical characterization, then to friction and wear tests using a pin-on-disc tribometer, at variable temperatures and humidity. Compared to C/C, they exhibited higher friction coefficients and wear rates, in line with the abrasive capacity of silicon carbide. For severe conditions, the friction coefficient sometimes dropped abruptly, through the occurrence of a mechanism close to lubrication. These results highlight both mechanical and physico-chemical effects of silicon carbide on the tribological properties of hybrid matrix composites at various energies.

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In aircraft construction, hybrid composites have become increasingly used as a skin material due to their beneficial characteristics in terms of fatigue and fire resistance. The material, consisting of alternate layers of aluminum and glass fiber-reinforced epoxy, is however complicated to model numerically. The work of Linde and DeBoer [109] deals with the modeling of hybrid composites and is focused on a detailed simulation of the inter-rivet buckling behavior in a stiffened fuselage shell. The skin material consists of a hybrid composite and is modeled with solid elements, layer by layer. Each of the constituents is provided with their respective material model. For the glass fiber-reinforced epoxy the user subroutine UMAT is employed for description of the failure modes, such as matrix cracking and fiber failure. The behavior of the delamination between the metal layer and the fiber-reinforced epoxy is also described with a user subroutine for interfaces, which is an optional contact definition. This subroutine contains a failure criterion for the delamination. The rivets between the stiffener and the skin of a fuselage are modeled by elastic solid elements with a plastic material model, but without any failure criterion. This reflects the design, where no rivet failure is allowed prior to buckling of the skin. A specially designed experimental test, which captures the main characteristics of inter-rivet buckling, is modeled and simulated. The numerical results are compared to the experimental data. Finally, conclusions and recommendations are given for future research. Polymer composite materials for aviation application based on carbon, organic and glass fibers and epoxy binders were exposed for 10 years under the warm humid climatic conditions of the city of Batumi on the Black Sea coast [110]. During this process, interlayer shear strength was measured and compared with other mechanical and physical characteristics measured in the center and surface layers of exposed plates. It has been shown that the surface of composites exposed to solar radiation is destroyed and disrupted to a greater extent than the inner layers. The gradient of mechanical properties across the thickness of a specimen can be significant. The investigation of polymeric composites related to a supersonic transportation application where polymer matrix composites utilized in primary structures were subjected to particular hygro thermal flight-cycles. In fact, the particular point in this study is the drying effect of supersonic flight at high temperature, around 130 ◦ C, on the durability of the material [111]. This phenomenon constitutes an entirely new situation for these materials in contrast with a classical subsonic flight at low temperature. The supersonic aircraft is supposed to be subjected to a succession of supersonic flight-cycles followed by a periodic maintenance operation. The objective of the study is first to characterize the in-service material state during the supersonic flight cycles and after the maintenance operations. Then, the challenge is to define the material geometry and environmental conditions to meet the in-service material state in short time. Thus, different accelerated cycles adapted to the new situation of supersonic flights, i.e. focusing on the cyclical drying of the material, are proposed. The applications of fiber-reinforced composites using polymer matrices have seen tremendous growth. In spite of the complexity of their behavior and the unconventional nature of fabrication and other aspects, the usage of such composites, even for primary load bearing structures in military fighters and transport aircraft, and satellites and space vehicles has been beneficially realized [112]. Most such usage constituted structural applications (such as in airframes) where service temperatures are not expected to he beyond 120 ◦ C. Attention is now focused on expanding the usage of such composites to other areas where temperatures could be higher – in the range 200–400 ◦ C. The intended applications are structural and nonstructural parts on or around the aero-engines and airframe components for supersonic or hypersonic aircraft. The development of polymer matrices – such as bis-maleimides, polyimides, cyanates, and liquid crystalline polymers and others – has brought such applications within the realm of practicability. The associated problems have been in terms of suitable processing technologies and in balancing the requirements of the performance with those of the processing. Ionic polymer-metal composites (IPMCs) are composites of a noble metal, conductive polymer or carbon/graphite, and charged polyelectrolyte membrane [113]. IPMCs have shown considerable progress in producing actuation in electric fields. These composites are also capable of sensing motion by producing a

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voltage difference when bent by a mechanical force. Work to date has yielded a force greater than 40 times the weight of an IPMC and large bending displacements with very low-input voltages. There is sufficient reason to believe that artificial muscles with viable strength can be produced with these composites. The IPMC, in addition to being resilient and elastic, is also lightweight and has a reaction speed that ranges from 1 microsecond to 1 second. For space missions, devices based on IPMCs will have numerous applications. On planetary surfaces, robotic arms and end effectors, motion-producing motors, actuators, and controllers are just a few examples of devices that can be produced using IPMCs. Effects of moisture into the adhesively bonded composite structures on their bond line mechanical strengths have been investigated [114–117]. The composite structures include a honeycomb sandwich structures fabricated by the co-cure and the pre-cure processes. The mass of moisture accumulated into the closed cells of the honeycomb sandwich panel specimens has been calculated. A pressure due to the vapor expansion in each cell of the sandwich panel has also been obtained for the minimum repair pressure to be applied to the laminate patched area by vapor pressure–temperature relations from the thermodynamic steam table, the ideal gas state equation and two vapor pressure equations. The bond line strengths of the laminated skins on the flat surface of honeycomb have been compared by the flat wise tension test and climbing drum peel test for dry, wet and repair-after-wet specimens, respectively. 23.3.2

Spacecraft Applications

Polymeric composite materials are fast gaining ground as the preferred materials in spacecraft construction. Their use as primary structural materials in recent years in several aerospace projects, e.g. in Space Shuttle Atlantis construction (Figure 23.19), has provided confidence leading to their acceptance as prime materials for aerospace vehicles. Although several applications in the aerospace sector are mentioned, the emphasis of this review is on applications of composites as structural materials where they have seen a significant growth in usage. Current solutions are briefly described and the scope for new developments is outlined. Polymer composites are widely used for high-temperature thermal protection (TPS) materials for spacecraft heat shields, and nozzle liners for solid propellant rocket motors [118]. Newly-developed silicon containing polymer, abbreviated MSP (Mitsui Silicon Containing Polymer), possesses high decomposition temperature

Figure 23.19

Space Shuttle Atlantis transported by a Boeing 747 Shuttle Carrier Aircraft.

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Figure 23.20

USSR rocket Soyuz.

(500 ◦ C), and high char yield (>94% at 1000 ◦ C) after pyrolysis. Therefore, the MSP polymer has potential as matrix resin of high-performance ablative composites compared with conventional phenolic resins. Thermal performance and ablation characteristics of a carbon composite with the MSP polymer (CF/MSP) were investigated under supersonic plasma air stream (e.g. in construction of USSR Rocket Soyuz depicted in Figure 23.20). The cold wall heat fluxes ranged from 1.2 to 11.0 MW/m2 . Conventional carbon fiber phenolic (CF/PR) composite was tested at the similar conditions for direct comparison. The mass loss of the CF/MSP composite was far less than that of the CF/PR composite, which is due to higher char yield and consequent high density of char layer. On the other hand, higher internal temperature was measured in the CF/MSP composite. In orbit, satellites are exposed to significant thermal variations. To ensure reliable operation of their onboard systems and equipment, thermal control of the spacecraft is necessary using cold, neutral or warm coatings. The Materials and Coatings Laboratory of the Thermal Control Services at CNES has elaborated a cold coating version by using a polysiloxane deposit on a metal substrate (such as polished aluminum or vacuum deposited silver). In geostationary orbit, polysiloxane, which has a high electrical resistivity (>1017 m), can accumulate implanted charges that can give rise to electrostatic discharges and damage the neighboring electronic systems. To prevent any electrostatic discharge problems in geostationary orbit, the resistivity of coatings should be reduced without altering their thermo-optical properties, in particular the low solar absorptivity and the high emissivity for cold coatings. Several methods have been studied, such as the incorporation of carbon nanotubes (CNT) and indium tin oxide (ITO) nanoparticles in the polysiloxane matrix, with the objective of attaining a high transparency, a high emissive, and an antistatic resin, which was tested in Space Shuttle Atlantis (Figure 23.21). The incorporation of CNT even with a lower content than the percolation threshold does not allow having good thermo-optical and electrostatic discharge (ESD) properties.

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Figure 23.21

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Atlantis taking off.

However, it is perhaps possible to diminish the percolation threshold by around one order of magnitude with a better dispersion of the CNT charges. In this case, it could be possible to obtain coatings with acceptable thermo-optical properties. Finally, the addition of the polysiloxane matrix of ITO nanoparticles allows to decrease the coating thickness and obtain very good thermo-optical and ESD properties [119]. Optical fiber sensors offer a number of advantages for spacecraft applications. A principal application is strain sensing for structural health monitoring, shape determination, and spacecraft qualification testing. Friebele et al. [120] reviewed the results of the work at the Naval Research Laboratory where optical fiber strain sensors have been used on spacecraft structures and ground test hardware. The sensors have been both surface mounted to the structure and embedded in fiber-reinforced polymer composites. The issue of potential strength reduction of high-performance composites due to embedded optical fiber sensors and leads has been studied, low-cost fabrication of tubular struts with embedded sensors has been demonstrated, and a novel technique for fiber ingress–egress from composite parts has been developed. Applications of fiber sensors include distributed dynamic strain monitoring of a honeycomb composite plate and a lightweight reflector during acoustic qualification tests, ultrahigh-sensitivity static strain and temperature measurements for precision structures, and on-line system identification of a lightweight laboratory truss. Structural polymer composite materials that can be converted into fuels and combusted with oxidizers for orbital propulsion of spacecraft have been developed [121]. The candidate materials have been identified and sustained combustion with nitrogen tetroxide (NTO) as an oxidizer has been demonstrated. To improve reaction chemistry, several energetic additives have been evaluated. Detailed material compatibility tests were conducted to identify stable combinations of structural polymer and energetic additives. The sustained combustion of structural polymeric materials with embedded additives and NTO has also been demonstrated. In the next phase of research, hydrogen peroxide as an oxidizer will be investigated. Samples of composites comprising thin metallic facesheets, structural polymer propellant matrix, and metallic mesh reinforcements (that also serve as electrical heaters/igniters for pyrolysis) were fabricated and their mechanical properties were measured. The concept of a spacecraft structural stringer, which also functions as a thruster, was developed using the composite material formulation. Both all-solid and hybrid stringer-thruster designs have been developed, and were used in the construction of the spacecraft Apollo 6 (Figure 23.22). All earth-orbiting spacecrafts are susceptible to impacts by meteoroids and pieces of orbital debris. These impacts occur at extremely high speeds and can damage flight-critical systems, which can in turn lead to

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Figure 23.22

Apollo 6 while dropping the interstage ring.

catastrophic spacecraft failure. With the advent of many new high-strength composite materials and their proliferation in aircraft applications, it is necessary to evaluate their potential for use in long-duration space and aerospace structural systems. Schonberg [122] presents the results of several experimental investigations in which several different composite materials were tested for their ability to prevent the perforation of the inner walls in multi-wall systems subjected to hypervelocity projectile impact. Damage to multi-wall systems that employ composite materials is compared to damage in traditional all-aluminum multi-wall structures of similar geometry and weight under similar impact conditions. The analyses performed show that the effectiveness of multi-wall systems with composite materials varies significantly depending upon the location within the multi-wall system of the composite materials. Based on the results obtained, it appears that composite materials are most effective when used as the middle wall of a triple-wall system. The results show that in double-bumper systems, Spectra/epoxy and Kevlar/epoxy inner bumpers were more effective in protecting the pressure wall against perforation by hypervelocity projectiles than equal-weight aluminum inner bumpers. The results obtained clearly demonstrate that using composite materials in combination with metallic materials such as aluminum in a multi-wall structure can significantly increase the protection afforded to the inhabitants of the structure over that provided by a traditional all-aluminum multi-bumper system of similar weight. A new technology is adopted to improve atomic oxygen resistance characteristics of spacecraft resin matrix composites [123, 124] which is to fill superfine particles into the resin matrix. The superfine plerospheres, which are non-reactive with atomic oxygen, are filled into the phenolic matrix in order to produce the AO-resistant fiber/phenolic resin composite in the paper. Plerospheres/fiber/phenolic resin composites were prepared and the AO exposure test was conducted in a ground-based AO effects simulation facility. It was found that the filling of plerospheres effectively protected the fiber/phenolic resin against AO erosion. After the AO exposure test of 50–60 h, the erosion yield of the glass fiber/phenolic resin and the carbon fiber/phenolic resin with plerosphere fillers could decrease to 9% and 13% of the composite without fillers, respectively. Furthermore, the erosion yield of all the composites decreased with the increase of the filler amount.

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23.4 Human Medicine Applications 23.4.1

Musculoskeletal and Bone Applications

The gradual shift from biostable prostheses to degradable, temporary implants represents one of the most significant trends in biocomposites application. In view of this trend, medical applications of degradable implant materials were reviewed with special emphasis on orthopedic polymeric implants. Various applications of different polymeric composites in human medicine are displayed in Figure 23.23 [125]. Among the polymeric implant materials derived from natural sources, collagen, various polysaccharides such as cellulose, and microbial polyesters have been intensively investigated. Among the synthetic, degradable polymers, aliphatic polyesters such as poly(glycolic acid), poly(lactic acid), poly(caprolactone) and polydioxanone are most commonly used. Only recently, several new classes of polymers such as poly (ortho esters), polyanhydrides, and degradable polycarbonates have been applied as implant materials. A particularly versatile group of new biomaterials with promising engineering properties are the ‘pseudo’- poly (amino acids), amino acid-derived polymers in which conventional peptide bonds have been replaced by various chemical linkages. Artificial legs are designed primarily to restore the ability to walk to amputees. The light weight, corrosion resistance, fatigue resistance, aesthetic appearance, and ease of fabrication of polymeric composite materials make them the ideal choice for modern limbs systems [125]. Thermoset polymer composites reinforced with glass, carbon or Kevlar fibers are widely used in these systems. A typical artificial leg system consists of three parts, namely socket, shaft, and foot. A prosthetic leg showing socket, shaft and foot is displayed in Figure 23.24. The new biodegradable starch-based polymeric composites have been applied in medicine are under consideration for use in orthopedic temporary applications and as tissue engineering scaffolds. It has been tested that by using these polymers it is possible to both produce polymer/hydroxyapatite composites (with or without the use of coupling agents) with mechanical properties matching those of the human bone, and to obtain three-dimensional structures generated by solid blowing agents that are suitable for tissue engineering applications [126]. The study focused on establishing the influence of several additives (ceramic fillers, blowing agents and coupling agents) and processing methods/conditions on the biocompatibility of the materials described above. The cytotoxicity of the materials was evaluated using cell culture methods, according to ISO/EN 109935 guidelines. A cell suspension of human osteosarcoma cells was also seeded on a blend of corn starch with ethylene vinyl alcohol and on corn starch with ethylene vinyl alcohol hydroxyapatite composites, in order to have a preliminary indication on cell adhesion and proliferation on the material’s surface. In general, the obtained results show that all the different materials based on corn starch with ethylene vinyl alcohol, (which are being investigated for use in several biomedical applications), as well as all the additives (including the novel coupling agents) and different processing methods required to obtain the different properties/products, can be used without inducing a cytotoxic behavior to the development of biomaterials. Composites of thermally-sensitive hydrogels and optically active nanoparticles have been developed for the purpose of photothermally modulated drug delivery [127]. Polymers of N-isopropylacrylamide (NIPAAm) and acrylamide (AAm) exhibit a lower critical solution temperature (LCST) that is slightly above body temperature. When the temperature of the copolymer exceeds the LCST, the hydrogel collapses, causing a burst release of any soluble material held within the hydrogen matrix. Gold–gold sulfide nanoshells, a new class of nanoparticles designed to strongly absorb near-infrared light, have been incorporated into poly(NIPAAmco-AAm) hydrogels for the purpose of initiating a temperature change with light; light at wavelengths between 800 and 1200 nm is transmitted through tissue with relatively little attenuation, absorbed by the nanoparticles, and converted to heat. Significantly enhanced drug release from composite hydro gels has been achieved in

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Dental Implant CF/C, SiC/C

Dental Post CF/C, CF/Epoxy, GF/Polyester

Arch Wire & Brackets GF/PC, GF/PP, GF/Nylon, GF/PMMA

Dental Bridges UHMWPE/PMMA CF/PMMA, GF/PMMA KF/PMMA

Dental Restorative Material Silica/BIS-GMA HA/2.2’(1-methacryloxydiethoxyphenyl)

Vascular Graft Cells/PTFE, Cells/PET PET/Collagen, PET/Gelatin PU/PU-PELA Abdominal Wall Prosthesis PET/PU, PET/Collagen

Intramedullary Nails CF/LCP, CF/PEEK GF/PEEK Tendon / Ligament PET/PHEMA, KF/PMA, KF/PE CF/ PTFE, CF/PLLA, GF/PU

Bone Replacement Material HA/PHB, HA/PEG-PHB CF/PTFE, PET/PU, HA/HDPE PET/PU, HA/PE, Bio-Glass/PE, Bio-Glass/PHB, Bio-Glass/PS, HA/PLA Spine Cage, Plate, Rods, Screws and Disc CF/PEEK, CF/Epoxy, CF/PS, Bio-Glass/PU, Bio-Glass/PS, PET/SR, PET/Hydrogel

Finger Joint PET/SR, CF/UHMWPE

Total Hip Replacement CF/Epoxy, CF/C, CF/PS, CF/PEEK, CF/PTFE, CF/UHMWPE, CF/PE, UHMWPE/UHMWPE

Bone Cement Bone particles/PMMA, Titanium/PMMA, UHMWPE/PMMA, GF/PMMA, CF/PMMA, KF/PMMA, PMMA/PMMA, Bio-Glass/Bis-GMA

Cartilage Replacement PET/PU, PTFE/PU, CF/PTFE CF/C Bone plate & Screws CF/PEEK, CF/Epoxy, CF/PMMA, CF/PP, CF/PS CF/PLLA, CF/PLA, KF/PC HA/PE, PLLA/PLDLA, PGA/PGA

Total Knee Replacement CF/UHMWPE UHMWPE/UHMWPE

External Fixation CF/ Epoxy

CF: carbon fibers, C: carbon, GF: glass fibers, KF: kevlar fibers, PMMA: Polymethylmethacrylate, PS: polysulfone, PP: Polypropylene, UHMWPE: ultra-high-molecular weight polyethylene, PLDLA: poly(L-DL-lactide), PLLA: poly (L-lactic acid), PGA: polglycolic acid, PC: polycarbonate, PEEK: polyetheretherketone: HA: hydroxyapatite, PMA: polymethylacrylate, BIS-GMA: bis-phenol A glycidyl methacrylate, PU: polyurethane, PTFE: polytetrafluoroethylene, PET: polyethyleneterephthalate, PEA: poltethylacrylate, SR: silicone rubber, PELA: Block co-polymer of lactic acid and polyethylene glycol, LCP: liquid crystalline polymer, PHB: polyhydroxybutyrate, PEG: polyethyleneglycol, PHEMA: poly(2-hydroxyethyl methacrylate)

Figure 23.23

Various applications of different polymeric composites in human medicine.

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Figure 23.24

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Prosthetic leg showing socket, shaft, and foot.

response to irradiation by light at 1064 nm. The release of methylene blue and proteins of varying molecular weight has been investigated. Additionally, the nano shell-composite hydrogels can release multiple bursts of protein in response to repeated near-IR irradiation. The predominant sensor materials used in medical acoustic imaging have traditionally been piezoelectric ceramics. More recently piezoelectric polymers, e.g. polyvinylidene fluoride, have been used in this area. Polymers present several advantages over ceramics, but have major drawbacks (poor emission constant, low coupling coefficient) [128]. Sershen et al. [128] investigated the electrical and mechanical properties of two different piezoelectric polymer composite materials (piezoflex1 and piezel). Practical transducer constructions with corona poled, PCB-backed and diced array architectures have been manufactured with the composites, and their performance evaluated in pulse echo mode. Poly-4-hydroxybutyrate (P4HB) is being developed as a new absorbable material for implantable medical applications [129]. P4HB promises to open up new opportunities for the development of medical applications by offering a new set of properties that are not currently available. The absorbable biomaterial is strong yet flexible, and degrades in vivo at least in part by a surface erosion process. While the chemical structure of P4HB is similar to that of current absorbable polyesters used in implantable medical products, P4HB is produced by a fermentation process rather than through a chemical synthesis. P4HB is a thermoplastic material that can be processed using standard plastics processing techniques, such as solution casting or melt extrusion. The strength of P4HB fibers prepared by melt extrusion compares well with that of traditional suturing materials; however, P4HB is typically more flexible. P4HB should find use in a wide variety of medical fields such as cardiovascular, wound healing, orthopedic, drug delivery, and tissue engineering applications. Natural bone is actually an inorganic/organic composite mainly made up of nanostructure hydroxyapatite (Ca10 (PO4 )6 (OH)2 , HAp), and collagen fibers. It is most important to synthesize nanocomposites of inorganic/organic in order to have good biocompatibility, high bioactivity and great bonding properties [130]. In this work, HAp nano-particle and HAp/chitosan (CTS) nanocomposite with a homogeneous microstructure

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were prepared and characterized. It is proposed that the nanostructure of hydroxyapatite/chitosan composite will have the best biomedical properties in the biomaterials applications. Recently, conducting polymers have attracted much interest in the development of biosensors [131]. The electrically conducting polymers are known to possess numerous features that allow them to act as excellent materials for immobilization of biomolecules and rapid electron transfer for the fabrication of efficient biosensors. In the present review [131] an attempt has been made to describe the salient features of conducting polymers and their wide applications in health care, food industries and environmental monitoring. Biodegradable synthetic polymers focusing on their potential in tissue engineering applications were applied. The major classes of polymers are briefly discussed with regard to synthesis, properties and biodegradability, and known degradation modes and products are indicated based on studies reported in the literature [132]. The vast majority of biodegradable polymers studied belongs to the polyester family, which includes polyglycolides and polylactides. Some disadvantages of these polymers in tissue engineering applications are their poor biocompatibility, release of acidic degradation products, poor processability and loss of mechanical properties very early during degradation. Other degradable polymers such as polyorthoesters, polyanhydrides, polyphosphazenes, and polyurethanes are also discussed and their advantages and disadvantages summarized. With advancements in tissue engineering it has become necessary to develop polymers that meet more demanding requirements. Recent work has focused on developing injectable polymer compositions based on poly (propylene fumarate) and poly (anhydrides) to meet these requirements in orthopedic tissue engineering. Polyurethanes have received recent attention in the development of degradable polymers because of their great potential in tailoring polymer structure to achieve mechanical properties and biodegradability to suit a variety of applications. Musculoskeletal tissue, bone and cartilage are under extensive investigation in tissue engineering research [133]. A number of biodegradable and bioresorbable materials, as well as scaffold designs, have been experimentally and/or clinically studied. Ideally, a scaffold should have the following characteristics: (i) three-dimensional and highly porous with an interconnected pore network for cell growth and flow transport of nutrients and metabolic waste; (ii) biocompatible and bioresorbable with a controllable degradation and resorption rate to match cell/tissue growth in vitro and/or in vivo; (iii) suitable surface chemistry for cell attachment, proliferation, and differentiation; and (iv) mechanical properties to match those of the tissues at the site of implantation. This paper [133] reviews research on the tissue engineering of bone and cartilage from the polymeric scaffold point of view. The wireless tadpole robot that has simple geometry, driven by low voltage and the undulatory fin-motion using IPMC (Ionic Polymer Metal Composite) actuator [134] have been investigated. The behavior of TadRob is tested and analyzed under various frequencies (1 ∼ 8 Hz) to find the correlation between actuator frequency and velocity of the robot. In addition, the robot velocity based on undulation motion and oscillation motion of the fin is compared to find the proper fin-motion in the viewpoint of. velocity efficiency for the robot. Also, steering capability is tested under variation of duty ratio. Based on experimental results, we can confirm that the velocity of TadRob can be controlled by changing frequency of input voltage and the steering angle can be increased by increasing the duty ratio. Composite materials, which can be very strong while having a low modulus of elasticity, are being studied because such materials have potential to be made into isoelastic hip prostheses. The injection molded CF/PEEK composite stem for total hip joint replacement is shown in Figure 23.25. Composites intended for medical applications incorporate carbon or polyamide as a fiber component, while polysulfone, polyetheretherketone, or polyethylene are used as a matrix component [135]. Mechanical properties (especially the modulus of elasticity) are emphasized because of the desire to match those properties of the proximal femur. Many of the variables that affect the mechanical properties of these materials are explained. The application of stress to different fiber orientations demonstrates the mechanical properties of the composite, and this is proved mathematically. It is shown that in composites with fibers oriented in the same direction, the modulus

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Figure 23.25

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Injection molded CF/PEEK composite for total hip joint replacement.

of elasticity in the direction of the fibers generally approaches that of the fibers as the amount of matrix decreases. Perpendicular to the fibers, the modulus of elasticity of the composite is only slightly greater than that of the matrix material. For isotropic chopped-fiber composites, the modulus of elasticity approaches that of the matrix as the fiber content decreases; at high-fiber content, the modulus is significantly less than that of oriented long-fiber composites. In general, the modulus of elasticity and fiber content have a linear relationship. Composites have fatigue properties that vary with direction and approach ultimate strength in tension but are lower in compression. The fatigue properties of proposed composites are discussed. Biodegradable polymers and bioactive ceramics are being combined in a variety of composite materials for tissue engineering scaffolds [136]. Materials and fabrication routes for three-dimensional scaffolds with interconnected high porosities suitable for bone tissue engineering are reviewed. Different polymer and ceramic compositions are applied and their impact on biodegradability and bioactivity of the scaffolds are discussed, including in vitro and in vivo assessments. The mechanical properties of today’s available porous scaffolds are analyzed in detail, revealing insufficient elastic stiffness and compressive strength compared to human bone. Further challenges in scaffold fabrication for tissue engineering such as biomolecules incorporation, surface functionalization and three-dimensional scaffold characterization are discussed, giving possible solution strategies. Stem cell incorporation into scaffolds as a future trend is briefly addressed, highlighting the immense potential for creating next-generation synthetic/living composite biomaterials that feature high adaptiveness to the biological environment. Because their chemical and physical properties may be tailored over a wide range of characteristics, the use of polymers is finding a permanent place in sophisticated electronic measuring devices such as sensors [137]. During the last five years, polymers have gained tremendous recognition in the field of artificial sensors in the goal of mimicking natural sense organs. Better selectivity and rapid measurements have been achieved by replacing classical sensor materials with polymers involving nanotechnology and exploiting either the intrinsic or extrinsic functions of polymers. Semiconductors, semiconducting metal oxides, solid electrolytes, ionic membranes, and organic semiconductors have been the classical materials for sensor devices. The developing role of polymers as gas sensors, pH sensors, ion-selective sensors, humidity sensors, biosensor devices, etc., are reviewed and discussed in this paper [137]. Both intrinsically conducting polymers and non-conducting polymers are used in sensor devices. Polymers used in sensor devices either participate in sensing mechanisms or immobilize the component responsible for sensing the analyte. Finally, current trends in sensor research and also challenges in future sensor research are discussed.

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Because of an aging population and increased occurrence of sports-related injuries, musculoskeletal disorders have become one of the major health concerns in the United States [138]. Current treatments, although fairly successful, do not provide the optimum therapy. These treatments typically rely on donor tissues obtained either from the patient or from another source. The former raises the issue of supply, whereas the latter poses the risk of rejection and disease transfer. This has prompted orthopedic surgeons and scientists to look for viable alternatives. In recent years, tissue engineering has gained increasing support as a method of treating orthopedic disorders. Because it uses principles of engineering, biology, and chemistry, tissue engineering may provide a more effective approach to the treatment of musculoskeletal disorders than traditional methods. This paper [138] presents a review of current methods and new tissue-engineering techniques for the treatment of disorders affecting bone, ligament, and cartilage. Biodegradable scaffolds have played an important role in a number of tissue engineering attempts over the past decade [139]. The goal of this review article is to provide a brief overview of some of the important issues related to scaffolds fabricated from synthetic biodegradable polymers. Various types of such materials are available; some are commercialized and others are still in the laboratories. The properties of the most common of these polymers are discussed here. A variety of fabrication techniques have been developed to fashion polymeric materials into porous scaffolds, and some of these are presented. The very important issue of scaffold architecture, including the topic of porosity and permeability, is discussed. Other areas, such as cell growth on scaffolds, surface modification, scaffold mechanics, and the release of growth factors, are also reviewed. A summary outlining the common themes in scaffold-related science, that are found in the literature, is presented. Functional graded materials (FGM) provide one new concept for guided tissue regeneration (GTR) membrane design with graded component and graded structure where one face of the membrane is porous, thereby allowing cell growth thereon, and the opposite face of the membrane is smooth, thereby inhibiting cell adhesion in periodontal therapy [140]. The goal of this study was to develop a three-layered graded membrane, with one face of 8% nanocarbonated hydroxyapatite/collagen/poly(lactic-co-glycolic acid) (nCHAC/PLGA) porous membrane, the opposite face of pure PLGA nonporous membrane, and the middle layer of 4% nCHAC/PLGA as the transition through layer-by-layer casting method. Then the three layers were combined with each other with flexibility and enough high mechanical strength as membrane because the three layers all contained PLGA polymer that can be easily used for practical medical application. The high biocompatibility and osteoconductivity of this biodegraded composite membrane was enhanced by the nCHAC addition, for the same component and nano-level crystal size with natural bone tissue. The osteoblastic MC3T3-E1 cells were cultured on the three-layered composite membrane; the primary result shows the positive response compared with pure PLGA membrane. Temperature-sensitive hyaluronic acid (HA) hydrogels were synthesized by photopolymerization of vinyl group modified HA in combination with acrylate group end-capped poly(ethylene glycol)–poly(propylene glycol)–poly(ethylene glycol) tri-block copolymer (Pluronic F127) [141]. The synthesized HA/Pluronic composite hydrogels gradually collapsed with increasing temperature over the range of 5–40 ◦ C, suggesting that the Pluronic component formed self-associating micelles in the hydrogel structure. Upon prolonged incubation in a buffer medium, the micelles slowly degraded due to the hydrolytic scission of the ester linkage between the Pluronic and acrylate group. The mass erosion occurred much faster at 37 ◦ C than at 13 ◦ C, indicating that at the higher temperature, the ester linkage between the Pluronic and acrylate group might be more exposed to an aqueous environment and thus be more readily hydrolyzed due to Pluronic micellization. Incorporation of recombinant human growth hormone in the hydrogel resulted in a sustained release profile which followed a mass erosion pattern. Ion-exchange Polymer Metal Composites (IPMC) is a new class of intelligent material that can be used effectively as actuators and artificial muscles. IPMC was fabricated and its displacement and force characteristics were investigated with respect to voltage, frequency and waveform of the controlling signal [142].

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A square waveform input generated slightly larger displacement and force than sinusoidal or triangular waveform. When the voltage was increased and the frequency was decreased, displacement and force were both increased. However, although the bending deformation of IPMC was large, the output force was much lower than expected. Improvement of the force output is key and is the main obstacle to be overcome in order to make IPMC of practical use. The number of recent findings in connection with ion-exchange polymer-noble metal composites (IPMC) as biomimetic sensors and actuators [143] have been reported. These smart composites exhibit characteristics of both actuators and sensors. Strips of these composites can undergo large bending and flapping displacement if an electric field is imposed across their thickness. Thus, in this sense they are large motion actuators. Conversely, by bending the composite strip, either quasi-statically or dynamically, a voltage is produced across the thickness of the strip between the two conducting electrodes attached. Thus, they are also large motion sensors. The output voltage can be calibrated for a standard size sensor and correlated to the applied loads or stresses. They can be manufactured and cut in any size and shape and in particular in the form of micro sensors and micro actuators for MEMS applications. In this paper [143] first the sensing capability of these materials is reported. The preliminary results show the existence of a linear relationship between the output voltage and the imposed displacement for almost all cases. Furthermore, the ability of these ionic polymer-metal composites as large motion actuators and robotic manipulators is presented. Several muscle configurations are constructed to demonstrate the capabilities of these IPMC actuators. This paper [143] further identifies key parameters involving the vibrational and resonance characteristics of sensors and actuators made with IPMCs. When the applied signal frequency is varied, so does the displacement up to a point where large deformations are observed at a critical frequency called resonant frequency where maximum deformation is observed. Beyond this, the actuator response is diminished. A data acquisition system was used to measure the parameters involved and record the results on a real time basis. Furthermore, reported in this paper [143] are load characterizations of such active polymer composites made with a noble metal such as platinum. The results show that these actuators exhibit good force-to-weight characteristics in the presence of low applied voltages. Finally reported are the cryogenic properties of these muscles for possible use by NASA in a harsh outer space environment of few Torrs and temperatures of the order of –140 ◦ C. These muscles are shown to work quite well in such a harsh cryogenic environment and thus present a great potential as sensors and actuators that can operate at cryogenic temperatures. The phenomenon of mass transfer-induced hydraulic actuation in ionic polymer-metal composites (IPMCs) strips with chemically plated surface electrodes was studied. A new category of actuators as ‘mass transfer induced hydraulic actuators (MTIHAs)’ is thus defined [144]. Such mass transport of ions and water is across the thickness of the strip and ends up near the permeable metal electrodes. The imposition of an electric field causes the mobile hydrated cations that are conjugated to the polymeric network anions to undergo electrophoretic dynamic migration that can result in local deformation of the material. Such electrophoretic behavior of the hydrated cations in IPMC causes the water to flow across the strip and partially leak out of the cathode side of the permeable electroded boundary. These leakages are undesirable because they lower the actuation performance. The IPMC is a good example of such MTIHAs. The advantage of MTIHAs is their potential to generate substantially high force densities if the leakage can be minimized or eliminated. Based upon this MTIHA concept, a certain manufacturing technique was developed to minimize the leakage and increase the force density of the conventional IPMCs by almost a factor of two (100% improvement in force). The trade-off design and fabrication of IPMC (ionic polymer metal composites) as an actuator for a flapping device have been described. Experiments for the internal solvent loss of IPMCs have been conducted for various combinations of cation and solvent in order to find out the best combination of cation and solvent for minimal solvent loss and higher actuation force [145]. From the experiments, it was found that IPMCs with heavy water as their solvent could operate longer. Relations between length/thickness and tip force of IPMCs

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were also quantitatively identified for the actuator design from the tip force measurement of 200, 400, 640, and 800 μm thick IPMCs. All IPMCs thicker than 200 μm were processed by casting NafionTM solution. The shorter and thicker IPMCs tended to generate higher actuation force but lower actuation displacement. To improve surface conductivity and to minimize solvent evaporation due to electrically heated electrodes, gold was sputtered on both surfaces of the cast IPMCs by the Physical Vapor Deposition (PVD) process. For amplification of a short IPMC’s small actuation displacement to a large flapping motion, a rack-and-pinion type hinge was used in the flapping device. An insect wing was attached to the IPMC flapping mechanism for its flapping test. In this test, the wing flapping device using the 800 μm thick IPMC could create around 10◦ ∼ 85◦ flapping angles and 0.5 ∼ 15 Hz flapping frequencies by applying 3 ∼ 4 V. A variety of bioactive composites have been investigated over the last two decades as substitute materials for diseased or damaged tissues in the human body [146]. In this paper, the rationale and strategy of developing these composites are given. Major factors influencing the production and performance of bioactive composites are discussed. Some promising composites for tissue replacement and regeneration are reviewed. On the basis of past experience and newly-gained knowledge, composite materials with tailored mechanical and biological performance can be manufactured and used to meet various clinical requirements. Organic–inorganic hybrid materials do not only represent a creative alternative to the design of new materials and compounds for academic research, but their improved and/or unusual features allow the development of innovative industrial applications [147]. Nowadays, most of the hybrid materials that have already entered the market are synthesized and processed by using conventional soft chemistry-based routes developed in the 1980s. These processes are based on: (a) the copolymerization of functional organosilanes, macromonomers, and metal alkoxides; (b) the encapsulation of organic components within sol–gel derived silica or metallic oxides; (c) the organic functionalization of nanofillers, nanoclays or other compounds with lamellar structures, etc. The chemical strategies (self-assembly, nanobuilding block approaches, hybrid MOF (metal organic frameworks), integrative synthesis, coupled processes, bio-inspired strategies, etc.) offered nowadays by academic research allow, through an intelligent tuned coding, the development of a new vectorial chemistry, able to direct the assembling of a large variety of structurally well-defined nano-objects into complex hybrid architectures hierarchically organized in terms of structure and functions. Looking to the future, there is no doubt that these new generations of hybrid materials, born from the very fruitful activities in this research field, will open a land of promising applications in many areas: optics, electronics, ionics, mechanics, energy, environment, biology, medicine (for example as membranes and separation devices), functional smart coatings, fuel and solar cells, catalysts, sensors, etc. Studies with biodegradable starch-based polymers have recently demonstrated that these materials have a range of properties, which make them suitable for use in several biomedical applications, ranging from bone plates and screws to drug delivery carriers and tissue engineering scaffolds [148]. The aim of this study was to screen the cytotoxicity and evaluate starch-based polymers and composites as potential biomaterials. The biocompatibility of two different blends of corn-starch, starch/ethylene vinyl alcohol (SEVA-C) and starch/cellulose acetate (SCA) and their respective composites with hydroxyapatite (HA) was assessed by cytotoxicity and cell adhesion tests. The MTT assay was performed with the extracts of the materials in order to evaluate the short-term effect of the degradation products. The cell morphology of L 929 mouse fibroblast cell line was also analyzed after direct contact with polymers and composites for different time periods and the number of cells adhered to the surface of the polymers was determined by quantification of the cryptozoic lactate dehydrogenate (LDH) activity. Both types of starch-based polymers exhibit a cytocompatibility that might allow for their use as biomaterials. SEVA-C blends were found to be the less cytotoxic for the tested cell line, although cells adhere better to SCA surface. The cytotoxicity test also revealed that SCA and SEVA-C composites have a similar response to the one obtained for SCA polymer. Scanning electron microscopy (SEM) analysis showed that cells were much more spread on the SCA polymer and LDH measurements showed a higher number of cells on this surface.

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Poly(L-lactic) acid (PLLA), polycaprolactone (PCL), three different copolymers based on poly(l-lactic) acid and polyglycolic acid (PLLA-co-PGA), and their composites with hydroxyapatite obtained from bovine bone (ossein), were tested in order to have information on the thermal, morphological, mechanical and biochemical properties in view of their use as biocompatible/biodegradable materials [149]. Ossein, which is essentially a biological hydroxyapatite, was found to improve the modulus and increase the hydrophilicity of the polymeric substrate. In addition, the size of the ossein particles was found to be critical for the improvement of mechanical properties. Finally, preliminary results on the in vitro biocompatibility of selected blends carried out by using primary cultures of human osteoblasts showed that the presence of hydroxyapatite stimulates a more positive cellular response. The preparation of bone tissue engineering is directly related to changes in materials technology [150]. While the inclusion of materials requirements is standard in the design process of engineered bone substitutes, it is also critical to incorporate clinical requirements in order to engineer a clinically relevant device. The CF/PEEK composite screws have been used for bones fixation (Figure 23.26). This review [150] presents the clinical need for bone tissue-engineered alternatives to the present materials used in bone grafting techniques, a status report on clinically-available bone tissue-engineering devices, and recent advances in biomaterials research. The discussion of ongoing research includes the current state of osseoactive factors and the delivery of these factors using bioceramics and absorbable biopolymers. Suggestions are also presented as to the desirable design features that would make an engineered device clinically effective. The development, fabrication and application of controlled porosity polymer-ceramic composite scaffolds, with 3-D interconnectivity, are designed to promote a richer supply of blood, oxygen and nutrients for healthy in-growth of bone cells [151]. Particulate-reinforced polymer-ceramic composites were developed by high shear mixing of polypropylene (PP) polymer and tricalcium phosphate (TCP) ceramic. Processing aids were used to improve plasticity and processibility to the composites. Controlled porosity scaffolds were fabricated via the fused deposition process, one of the commercially-available rapid prototyping (RP) techniques. These porous scaffolds were characterized for their use as bone grafts in terms of physical, mechanical and biological properties. Hg-porosimetry was performed to determine the pores size and distribution. Scaffolds with different complex internal architectures were also fabricated using this composite material. Tensile

Figure 23.26

CF/PEEK composite screws.

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properties of neat PP (as received), PP with processing aids (without TCP) and PP-TCP composite (with processing aids) were evaluated and compared using standard dog bone samples. Uniaxial compression tests were performed on cylindrical porous samples with an average pore size of 160 μm and varying vol% porosity (36%, 48% and 52%). Samples with 36 vol% porosity showed the best compressive strength of 12.7 MPa. Cytotoxicity and cell proliferation studies were conducted with a modified human osteoblast cell-line (HOB). Results showed that these samples were nontoxic with excellent cell growth during the first two weeks of in vitro testing. Hydroxyapatite (HA)-reinforced polymers have been proposed as a method of improving the biological properties of bone cements and implant materials. For example, bone cements based on polymethylmethacrylate (PMMA) have long been used to secure orthopedic implants to the skeleton [152]. This composite could also be used as a polished coating on other materials or in bulk form, shaped or molded, to custom fit a specific clinical need. However, complications may occur as a result of the limited mechanical and biological properties of PMMA. The purpose of this investigation was to determine whether the incorporation of HA in a PMMA matrix would enhance the biological properties of osteoblast response as compared to PMMA alone. Fetal rat calvarial osteoblasts were plated on discs of PMMA, PMMA/HA, commercially pure titanium (CpTi) and tissue culture polystyrene (control). Osteoblast attachment and day 2 proliferation were similar on all implant materials, whereas day 8 proliferation on PMMA/HA was significantly higher than on PMMA and similar to CpTi and control. Extracellular matrix production was examined by immunohistochemistry which indicated that osteoblasts cultured on PMMA/HA showed a more distinct networked pattern of organized fibronectin. Histochemical staining of mineralization was examined by confocal microscopy, which demonstrated a higher degree of mineralization in nodules formed on PMMA/HA as compared to PMMA. Together, these results indicate that the addition of HA in a PMMA matrix improves osteoblast response as compared to PMMA alone. Therefore, the incorporation of HA into a PMMA matrix may be a useful method to provide PMMA materials with enhanced osteogenic properties. Bioactive composites comprising synthetic hydroxyapatite (HA) particulate and semi-crystalline polyetherketone (PEEK) polymer were produced for biomedical application [153]. The bone plates made of CF/epoxy and CF/PEEK composite material are shown in Figure 23.27. Particulates were incorporated into a PEEK polymer matrix through a series of processing stages involving melt compounding, granulating and injection molding. This investigation presents the processing route employed and the mechanical properties of HA–PEEK composites. In general, Young’s modulus, compressive strength and micro-indentation hardness increased with increasing amount of HA particulate.

(a)

(b)

Figure 23.27

Bone plates made of CF/epoxy (a) and CF/PEEK (b) composite material.

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On the other hand, tensile strength and strain to failure decreased with increasing HA loading. The tensile strength and Young’s modulus of HA–PEEK composites were found to be within the bounds of bony tissue. These results suggest that the bioactive HA–PEEK composites have the potential for use as an alternative material for load-bearing orthopedic application. The fabrication process for a novel carbon fiber-reinforced polymer (polyamide 12) composite femoral stem using inflatable bladder molding was studied. The effects of processing temperature, holding time and applied internal pressure on the consolidation quality of the composite were investigated [154]. Consolidation quality was evaluated by density and void content measurements and scanning electron microscope analysis. As expected, void content (porosities) and presence of large resin pockets were found to increase for lower processing temperature, holding time and applied pressure. Crystallinity as well as melting temperatures measured using differential scanning calorimetry could be related to molding conditions. A progressive reduction of the previous thermal history (crystalline peak of neat composite) and an increase in crystallinity were obtained for higher molding temperature. Static compression testing with void content analysis of molded specimens was used to determine optimal molding conditions. The composite structure molded showed compressive modulus close to cortical bones. Compression load at failure of composites molded in optimal conditions were found to be three times higher than those of femoral bone for jumping on one leg or 10 times those for normal gait. The molded composite structure appears to be an excellent candidate for femoral stems used in total hip arthroplast as well as for various medical and prosthetic components production (Figure 23.28). During the past few years new composite materials for nerve regeneration have been applied [155, 156], and new polymeric composites for bone and cartilage repair and long-term implants have been tested [157–160].

23.4.2

Dentistry Applications

The currently-used commercial restorative composites contain a mixture of various crosslinking dimethacrylates, glass- and/or silicon dioxide fillers, and a photoinitiator system. They are cured by irradiation with

Figure 23.28

Medical and prosthetic components.

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visible light [161]. New developments of polymeric composites for restorative filling materials are mainly focused on the reduction of polymerization shrinkage, and improvement of biocompatibility, wear resistance and processing properties. This can be partially achieved by using new tailor-made monomers and optimized filler particles. The contribution [161] describes the polymeric chemical aspects of the application of new monomers, e.g. cyclic monomers, liquid-crystalline monomers, ormocers, branched monomers, compomers or Bis-GMA analogues or substitutes for restorative composites. In addition, the contribution of new filler technologies for the improvement of restorative composites is discussed. The placement of dental composites is complicated by the contraction that accompanies polymerization of these materials [162]. The resulting shrinkage stress that develops during cure of a bonded restoration can induce defects within the composite, the tooth or at the interface resulting in compromized clinical performance and/or esthetics. In light of the substantial efforts devoted to understanding and attempting to control shrinkage stress and strain in dental composite restoratives, this paper [162] offers a perspective on the conversion-dependent development of shrinkage and stress. The relationships between polymer property development and the physical evolution of the network structures associated with dental polymers as well as the interrelated kinetics of the photo polymerization reaction process are examined here. Some of the methods used to assess conversion in dental resins and composites are considered. In particular, newlyintroduced techniques that allow real-time analysis of conversion by near-infrared spectroscopy to be coupled directly to simultaneous dynamic measurements of either shrinkage stress or strain are described. The results are compared with reports from the dental materials literature as well as complementary studies in other related fields of polymer science. The complex, nonlinear correlation between conversion, shrinkage and stress are highlighted. The development and continued evolution of photopolymerizable dental materials, particularly dental composite restoratives, represents a significant practical advance for dentistry. The finished dental composite prosthesis in displayed in Figure 23.29 [163]. The highly-successful integration of the light-activated curing process for dental applications is described in this review. The basic mechanisms by which the photoinitiators efficiently convert monomers into polymers are discussed, along with the variety of factors that influence the photopolymerization process. The conventional camphorquinone-amine visible light photoinitiator system used in most dental restorative materials is illustrated as well as some alternative initiator systems that have been studied for dental materials applications. The application of fibrous composite polymeric materials in dentistry and orthopedics was reviewed [164]. Furthermore, the authors introduced functionally-graded composite dental post, aesthetic composite arch wires and brackets, and braided carbon/PEEK composite compression bone plate. Functionally-graded composite

Figure 23.29

Finished dental composite prosthesis – lateral view.

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dental post has continuously varied stiffness and this function successfully solved stress concentration at the root. Aesthetic composite arch wires made of glass/epoxy unidirectional composite have been targeted to obtain the highest bending performance by interface control. An epoxysilane coupling agent with 1.0 wt% solvent showed the highest bending performance. Aesthetic composite bracket fabricated by braided preform in order to reinforce tie-wing slot where orthodontic force is applied through an arch wire. Composite brackets have indicated around 43% mechanical resistance of stainless-steel bracket at the tie-wing slot. The dental composites were based on BTDMA and a new dimethacrylate monomer was based on BTDA (3,3 ,4,4 -benzophenone tetra carboxylic dianhydride). The polymeric composites were prepared by mixing the silane-treated filler with the monomers [165]. The prepared pastes were inserted into the test molds and heat cured. Light-cured composites were also prepared using camphorquinone and amine as photoinitiator system. The results showed that the mechanical properties of the new composite are comparable with the properties of the Bis-GMA-based composite but its water sorption is higher. BTDMA as a monomer containing aromatic rings and carboxylic acid groups in its structure gives a composite with good mechanical properties. There is a close relation between the contact angle, water sorption of the cured composite and water uptake of their monomers. Finding new monomers as alternatives for Bis-GMA have been a challenge in the field of dental materials and any investigation into the properties of new composites would be beneficial in the development of dental materials. Short fiber-reinforced polymer composites are nowadays used in numerous tribological applications, e.g. for dental application. In spite of this fact, new developments are still under way to explore other fields of application for these materials and to tailor their properties for more extreme loading conditions [166]. The references given at the end of this review [166] describe some of these developments. In the present overview further approaches in designing polymeric composites in order to operate under low friction and low wear against steel counterparts are described. A particular emphasis is focused on special filler- (including nanoparticle) reinforced thermoplastics and thermosets. Especially, the influence of particle size and filler contents on the wear performance is summarized. In some of the cases, an integration of traditional fillers with inorganic nanoparticles is introduced and presents an optimal effect. Furthermore, some new steps towards the development of functionally-graded tribo-materials are illustrated.

23.5 Electrical and Electronic Applications Conducting polymeric composites have found two main kinds of application in electronics so far: as materials for construction of various devices and as selective layers in chemical sensors [167]. In either case, interaction with ambient gases is critical. It may compromise the performance of a device based on conducting polymers, whereas it is beneficial in a sensor. Conductivity has been the primary property of interest. The aim here is to discuss the usability of conducting polymeric composites in both types of electronic applications in light of these two parameters. Printed circuit boards are big consumers of polymeric composite materials. PCB are compression molded to form rigid substrates that support the silicon chips and interconnect the other electrical components in the wide range of electronic devices. The majority of PCB are made with E-glass/epoxy prepreg, although other reinforcing fibers, including aramid and quartz, are sometimes used for specialty applications. The printed circuit boards based on epoxy/glass composites are displayed in Figure 23.30 [168]. Resin alternatives include vinyl ester and polyester for commodity boards, and cyanate esters, polyimide and bismaleimide triazine for more demanding, elevated-temperature applications. PCB design complexity has increased as electronic devices have become smaller and endowed with more functions, which is pushing suppliers to create thinner laminates for higher layer counts and greater circuit density. Nonwoven aramid fiber mat, wet out with epoxy

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Figure 23.30

Printed circuit boards based on epoxy/glass composites.

or polyimide, is an alternative laminate substrate material for high-performance applications, such as avionic and ultra-thin cell pones. Embedded capacitor technology can increase silicon packing efficiency, improve electrical performance, and reduce assembly cost compared with traditional discrete capacitor technology [169]. Developing a suitable material that satisfies electrical, reliability, and processing requirements is one of the major challenges of incorporating capacitors into a printed wiring board (PWB). Polymer–ceramic composites have been of great interest as embedded capacitor material because they combine the processability of polymers with the high dielectric constant of ceramics. A novel nanostructure polymer–ceramic composite with a very high dielectric constant (εr ∼ 110, a new record for the highest reported εr value of a nanocomposite) was developed in this work. A high dielectric constant is obtained by increasing the dielectric constant of the epoxy matrix (εr >6) and using the combination of lead magnesium niobate-lead titanate (PMN-PT)/BaTiO3 as the ceramic filler. This nanocomposite has a low curing temperature (<200 ◦ C); thus, it is multichip-module laminate (MCM-L) process-compatible. An embedded capacitor prototype with a capacitance density of 50 nF/cm2 was manufactured using this nanocomposite and spin-coating technology. The effect of the composite microstructure on the effective dielectric constant was studied. This novel nanocomposite is used for integral capacitors in PWBs. Carbon nanotube–polymer composites have been researched extensively for many applications requiring the combination of unique electronic, optical, and/or mechanical properties of carbon nanotubes and polymer materials [170]. However, to build integrated carbon nanotube–polymer-based systems, it is necessary to have a stateof-the art facility to incorporate organized nanotube architectures in selected polymer matrixes as well as engineer the interfaces between the two constituents. Here is presented a direct and effective method for fabricating flexible carbon nanotube-polymer composites by incorporating aligned and patterned multi-walled carbon nanotubes (MWNT) into a soft poly(dimethylsiloxane) (PDMS) matrix. The fabrication and electrical characterization of a flexible hybrid composite structure using aligned multi-wall carbon nanotube arrays in a poly(dimethylsiloxane) (PDMS) matrix was realized. Using lithographically patterned nanotube arrays, one can make these structures at any length scale from submicrometer levels to bulk quantities. The PDMS matrix undergoes excellent conformal filling within the dense nanotube network, giving rise to extremely flexible conducting structures with unique electromechanical properties. Its robustness is demonstrated against highstress conditions, under which the composite is found to retain its conducting nature. It is also demonstrated that these structures can be utilized directly as flexible field-emission devices. The devices show some of the best field-enhancement factors and turn-on electric fields reported so far.

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Recently, conducting polymers have attracted much interest in the development of biosensors [171]. The electrically conducting polymers are known to possess numerous features that allow them to act as excellent materials for immobilization of biomolecules and rapid electron transfer for the fabrication of efficient biosensors. In the present review [171], an attempt has been made to describe the salient features of conducting polymers and their wide applications in health care, food industries, environmental monitoring, etc. Conducting polyaniline forms an important family of electronic polymers with a developed potential application for a number of areas because of its flexible chemistry, processibility, environmental stability and ease of forming composites [172]. The electromagnetic interference shielding effectiveness of conducting polyaniline (PANI)–ABS composites was studied at 101 GHz. It was observed that shielding effectiveness of the PANI–ABS composites increases with the increase in the loading levels of the conducting polymer doped with hybrid dopants. The lower loading of PANI doped with hybrid dopants in the molded conducting composites can be effectively used for the dissipation of electrostatic charge. However, with higher loadings, a shielding effectiveness of 60 dB has been achieved which makes the conducting composites a potential EMI shielding material for its application in encapsulation of electronic equipments in electronic and in high-tech applications. The developments of composite membranes for polymer electrolyte fuel cell (PEFC) applications has been discussed [173]. The limitations of per-fluorinated polymer electrolyte membranes to low temperature (<80 ◦ C) PEFC applications was studied. Research on alternative proton-conducting membranes to the per-fluorinated membranes for high-temperature PEFC applications are shown. The development of the bis[(perfluoroalkyl)sulfonyl]imide as an alternative membrane to the per-fluorinated family is indicated. The concept of synergetic composite membranes for high temperature PEFC applications is introduced. Recent approaches and concepts for the elaboration of new composites membranes are described. The following aspects of the researches on proton conducting proton membranes are discussed: (a) macro- and nanocomposites per-fluorinated ionomer composite membranes (PFICMs); (b) partially per-fluorinated composite membranes; (c) non-per-fluorinated composite membranes. Results based on original works are also presented. In each case, the type of composite membrane is well described. Accordingly, organic–inorganic, organic–organic, organic–acid and organic–base complexes composite membranes are considered. The challenges related to these developments are discussed. Prospective applications of composite membranes for high-temperature PEFC applications are discussed. Porous polytetrafluoroethylene (PTFE) membranes were used as support material for Nation/PTFE composite membranes. The composite membranes were synthesized by impregnating porous PTFE membranes with a self-made Nafion solution [174]. The resulting composite membranes were mechanically durable and quite thin relative to traditional perfluorosulfonated ionomer membranes (PFSI); the composite membranes are expected to be of low resistance and cost. In this study [174], three kinds of porous PTFE films were used to prepare Nafion/PTFE composite membranes of different thickness. Scanning electron micrographs and oxygen permeabilities showed that Nafion resin is distributed uniformly in the composite membrane and completely plug the micropores; there is a continuous thin Nafion film present on the PTFE surface. The variation in water content of the composite and Nafion 115 membranes with temperature was determined. At the same temperature, water content of the composite membranes was smaller than that of the Nafion 115. In both dry and wet conditions, maximum strength and break strength of C-325# and C-345# were larger than those of Nafion 112 due to the reinforcing effect of the porous PTFE films. And the PEMFC performances and the lifetime of the composite membranes were also tested on the self-made apparatus. Results showed that the bigger the porosity of the substrate PTFE films, the better the fuel cell performance; the fuel cell

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performances of the thin composite membranes were superior to that of Nafion 115 membrane; and after 180 h stability test at 500 mA/cm2 , the cell voltage showed no obvious drop. The performance of supercapacitors with multi-walled carbon nanotubes deposited with conducting polymer as activate materials was greatly enhanced in contrast with electric double-layer supercapacitor with carbon nanotubes due to the conducting polymer’s faradaic effect. They are promising as the secondary power sources in electric vehicles propulsion [175]. Polypyrrole and poly(3-methylthiophene) were uniformly deposited onto multi-walled carbon nanotubes in organic system by chemical methods. A carbon nanotubespolypyrrole composite–carbon nanotubes-poly(3-methyl-thiophene) composite-based supercapacitor prototype (CNTs-pPy–CNTs-pMeT SCP), a carbon nanotubes–carbon nanotubes-polypyrrole based hybrid SCP (CNTs–CNTs-pPy SCP), a carbon nanotubes–carbon nanotubes-poly(3-methylthiophene)-based hybrid SCP (CNTs–CNTs-pMeT SCP) as well as a CNTs–CNTs corresponding SCP were assembled in 1 M LiClO4 acetonitrile solution. Their voltammetry characteristics, galvanostatic discharge and AC. impedance spectra were carried out in two-electrode mode. Pseudo capacitance effects are found out from those SCPs with composite electrodes and their measured capacitances are 87, 45, and 72 F.g−1 for CNTs-pPy–CNTs-pMeT SCP, CNTs–CNTs-pMeT SCP and CNTs–CNTs-pPy SCP, respectively. They are much larger than that of 21 F.g−1 for the CNTs–CNTs corresponding SCP, which is a double-layer SCP. Their measured specific energy is 1.82, 0.88 and 1.33 W.h.kg−1 for those SCPs with composite electrodes. They are also much higher than that of 0.58 W h.kg−1 for the CNTs–CNTs corresponding SCP. Because their chemical and physical properties may be tailored over a wide range of characteristics, the use of polymers is finding a permanent place in sophisticated electronic measuring devices such as sensors [176]. During the last five years, polymers have gained tremendous recognition in the field of artificial sensors with the goal of mimicking natural sense organs. Better selectivity and rapid measurements have been achieved by replacing classical sensor materials with polymers involving nanotechnology and exploiting either the intrinsic or extrinsic functions of polymers. Semiconductors, semiconducting metal oxides, solid electrolytes, ionic membranes, and organic semiconductors have been the classical materials for sensor devices. The developing role of polymers as gas sensors, pH sensors, ion-selective sensors, humidity sensors, biosensor devices, etc., are reviewed and discussed in this paper [176]. Both intrinsically conducting polymers and nonconducting polymers are used in sensor devices. Polymers used in sensor devices either participate in sensing mechanisms or immobilize the component responsible for sensing the analyte. Finally, current trends in sensor research and also challenges in future sensor research are discussed. Composites of high molecular weight polyaniline (PANI) and various weight percentages of singlewalled carbon nanotubes (SWCNT) were fabricated using solution processing [177]. Electrical characteristics of metal semiconductor (MS) devices fabricated from the PANI/SWNT composites were studied. Current–voltage (I–V) characteristics of these devices indicate a significant increase in current with an increase in carbon nanotube concentration in the composite. The dominant transport mechanisms operating in these devices were investigated by plotting the forward I–V data on a log–log scale, which revealed two power-law regions with different exponents. In the lower voltage range, the exponent is approximately 1, implying that the charge transport mechanism is governed by Ohm’s law. The charge transport mechanism in the higher voltage range, where the exponent varies between 1.1 and 1.7, is consistent with space-charge-limited (SCL) emission in the presence of shallow traps. The critical voltage (V c ), which characterizes the onset of SCL conduction, decreases with increasing SWNT concentration. In addition, V c was observed to increase with temperature. These initial results indicate that with further improvements in material consistency and reduction in defect densities, the polyaniline/single-walled carbon nanotube composite material can be used to fabricate organic electronic devices leading to many useful applications in microelectronics. A Schottky barrier diode based on a composite of polyaniline with polystyrene has been fabricated and characterized using indium as the Schottky contact and platinum as an ohmic contact [178]. Current–voltage (I–V) plots were nonlinear and capacitance–voltage (C–V) plots were almost linear in reverse bias indicating

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rectification behavior. Various junction parameters were calculated from the temperature dependent I–V and C–V data and discussed. These results indicate that the composite materials have better mechanical strength and diode quality compared to that of pure polymer. An approach to fabricate ferromagnetic/polymer composite nanotubes has been developed [179]. The surfaces of the pores in self-ordered porous alumina membranes are wetted with a polystyrene or poly-llactide layer containing a metallo-organic precursor. Decomposition of the precursor leads to the formation of thin-walled magnetic tubes with diameters of 160–450 nm and wall thicknesses of a few nanometers. The magnetic properties of the tube arrays are interpreted as a result of the tube morphology and microstructure. The novel type of rechargeable lithium polymer battery based on the combination of a nanocomposite electrolyte and a cathode of the phospho-olivine family [180] has been prepared. The results demonstrate features, in terms of power capability and cyclability, largely exceeding those so far reported for conventional lithium polymer batteries. These unique performances make this battery suitable for applications in electric vehicles. The bio-based composite material, suitable for electronic as well as aeronautical and automotive applications, was developed from soybean oils and keratin feather fibers (KF) [181]. This environmentally- friendly, low-cost composite can be a substitute for petroleum-based composite materials. Keratin fibers are a hollow, light, and tough material and are compatible with several soybean (S) resins, such as acrylated epoxidized soybean oil (AESO). The new KFS lightweight composites have a density ρ≈1 g/cm3 , when the KF volume fraction is 30%. The hollow keratin fibers were not filled by resin infusion and the composite retained a significant volume of air in the hollow structure of the fibers. Due to the retained air, the dielectric constant, k, of the composite material was in the range of 1.7–2.7, depending on the fiber volume fraction, and these values are significantly lower than the conventional silicon dioxide or epoxy, or polymer dielectric insulators. The coefficient of thermal expansion (CTE) of the 30 wt% composite was 67.4 ppm/◦ C; this value is low enough for electronic application and similar to the value of silicon materials or polyimides used in printed circuit boards. The water absorption of the AESO polymer was 0.5 wt% at equilibrium and the diffusion coefficient in the KFS composites was dependent on the keratin fiber content. The incorporation of keratin fibers in the soy oil polymer enhanced the mechanical properties such as storage modulus, fracture toughness, and flexural properties, ca. 100% increase at 30 vol%. The fracture energy of a single keratin fiber in the composite was determined to be about 3 kJ/m2 with a fracture stress of about 100–200 MPa. Considerable improvements in the KFS composite properties should be possible by optimization of the resin structure and fiber selection. Functional macromolecules once sought almost exclusively for their mechanical or thermal properties are now finding numerous applications in a variety of areas. In some cases, newly-designed functional polymers provide such an array of new properties and functions that they effectively create fields rather than extend or support them. This brief account [182], focusing largely on the work of the author’s own laboratory, explores recent advances in the design of functional polymers in two seemingly very disparate areas: organic or ‘plastic’ electronics and polymer therapeutics. While at first glance the requirements for functional polymers appear to be quite different for these two areas, the same general design concepts apply, focusing first on the definition of the key properties to be achieved, followed by the establishment of a synthetic blueprint encompassing these aims, and concluding with proof of concept studies. The examples given in this account [182] illustrate the vibrant nature of research in the area of functional polymers and its great potential impact on tomorrow’s technologies. The layer-by-layer (LbL) nano-assembly for deposition of ultrathin poly(anilinesulfonic acid) (SPANI) films for fabrication of highly-sensitive and rapid response humidity sensors has been used. Spin coating was also used for fabrication of SPANI-based humidity sensors for comparison [183]. The change in electrical sheet resistance of the sensing film was monitored as the device was exposed to humidity. For a 5% change in relative humidity, the sensitivity was measured to be 11 and 6% from layer-by-layer based and the

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spin-coated humidity sensors, respectively. An intended application for these layer-by-layer assembled devices is in disposable handheld instruments to monitor the presence of humidity in humidity-sensitive environments. Conducting polyaniline/tungsten oxide (WO3 ) composites have been synthesized by an in situ deposition technique by placing fine graded WO3 in a polymerization mixture of aniline [184]. This is a single step polymerization process for the direct synthesis of emeraldine salt phase of the polymer. The results were also well supported by FTIR spectral analysis, scanning electron microscope (SEM), XRD and conductivity measurements. High temperature conductivity measurements show thermal activated behavior. The change in resistance with respect to percent relative humidity (RH) is observed. The composites in the pellet form exhibit almost linear behavior within a chosen range of humidity (ranging between 10 and 95 % RH). An electronic nose that uses an array of 32 polymer–carbon black composite sensors has been developed, trained, and tested [185]. By selecting a variety of chemical functionalities in the polymers used to make sensors, it is possible to construct an array capable of identifying and quantifying a broad range of target compounds, such as alcohols and aromatics, and distinguishing isomers and enantiomers (mirror-image isomers). In many applications, polymer-supported metal films as interconnects for large area application in electronics have been studied [186–189]. A shadow-masked electro spray method was used to fabricate a highperformance carbon black–poly(vinyl pyrrolidone) (CB–PVP) composite sensor on a position-selected area of a sensor substrate [190]. This new approach, when compared to the common drop-casting method, was shown to be very advantageous for the preparation of an active layer with optimum average thickness and porous microstructure, which is important to obtain high sensitivity and fast detection times. The thickness obtained was proportional to the electro spray scan number, and the field-assisted generation of small droplets led to the formation of rough, porous microstructures. The CB content-dependent variation in sensor resistance exhibited percolation behavior with an abrupt decrease in the specific middle region with increasing CB content. The most sensitive methanol detection was observed for sensors with slightly larger CB content rather than that of the most rapidly varying midpoint in a percolation curve, which was probably due to increased susceptibility to noise resulting from high porosity. Polymer concentration was also observed to have a significant effect on sensing properties owing to a change in film morphology; the optimum concentration was approximately 0.25%, at which the conditions were suitable for the formation of a porous, sensitive sensor with a reasonable deposition speed.

23.6 Conclusion Advancements in polymeric composites technologies have been the driving force over the years in all applications. The important properties of polymeric composites, i.e. light weight and high strength, have made them a natural substitute for multiple components for automotive, naval, aircraft, and spacecraft construction. The carbon-based polymeric composites represent the dream of automotive engineers due to the light, strong, stiff and stable properties of these materials. The acceleration increase of the vehicles can be reached by reducing the mass of the construction material, and among mass-reducing materials polymeric composites rank highly. The polymeric composites used for prosthesis construction must have some special properties, namely long-term durability and reliability. The further development of polymeric composites can be supported by research into long-term durability in the human body. The application of polymeric composites in human medicine is very attractive due to the possibility of reaching the performance properties, which are comparable with original human tissues. The application of polymer composites in microelectronics has the potential to open new directions in integration and performance of devices and personal computer construction. Conformal form factors to provide the ability to install the large area electronics in a variety of locations. Even more difficult is to not just integrate a flexible substrate, but also to improve electrical

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performance. The keys to achieving the desired levels of functionality are advances in device technology that can be manufactured cost effectively over large areas and process methods, but to achieve the device/circuit performance for demanding applications, significant improvement in materials and device characteristics must be reached.

References 1. P. K. Mallick, Fiber reinforced composites (2nd ed.), Marcel Dekker, New York, 1993. 2. P. K. Mallick, Introduction: Definitions, Classifications, and Applications in Composites Engineering Handbook (P.K. Mallick, Ed.), Marcel Dekker, New York, 1997, p. 1–50. 3. S.K. Mazumdar, Composites Manufacturing: Materials, Product, and Process Engineering, CRC Press, New York, 2002. 4. S. Black, “The World of Composites: An Overview,” in Sourcebook 2007, The International Guide to Products and Services for the Composites Industry, J. Sloan, Ed., Ray Publishing, 2007. 5. D. Mattsson, Mechanical Performance of NCF Composites, Ph.D. Thesis, Lule˚a University of Technology, Department of Applied Physics and Mechanical Engineering (2005). 6. B. Z. Jang, Advanced polymer composites, ASM International, New York, 1994. 7. “Materials for Lightweight Vehicles”, Office of Transportation Materials, U.S. Department of Energy, Washington D.C., July 1992. 8. H. Baker, Liquid Molding Technology Rolls Ahead, Adv. Mater. Proc., Feb. 1991, p. 39–34. 9. PPO Appearance Part is Blow Molded, Modern Plast., Vol 67, Feb. 1990, p. 19–20. 10. S.A. Wood, In Subtle Ways, Plastic Tighten an Alliance with the Automotive Market, Modern Plast., April 1991, p. 46–49. 11. L. Dodyk, High Performance Composites for the Automotive Industry, Adv. Comp., June 1991, p. 21–27. 12. D.F. Baxter, Jr., Green Light to Plastic Engine Parts, Adv. Mater. Proc., May 1991, p. 26–31. 13. P. Beardmore, J.J. Harwood, and E.J. Horton, Design and Manufacture of a GrFRP Concept Automobile, Proc. Intl. Conf. Compos. Materials, Elsevier, Paris, Aug. 1980, p. 47–60. 14. H.T. Kulkarni and P. Beardmore, Design Methodology for Automotive Components Using Continuous Fiber Reinforced Materials, Composites, Vol 12, 1980, p. 225–235. 15. D. Vesey and S. Abouzadbr, E-Coat-Capable Plastics – A New Generation of Materials for Automotive Exterior Body Panels, in Automotive Exterior Body Panels, Society of Automotive Engineers, Detroit, 1988, p. 1–6. 16. J. Arimond and B.B. Fitts, Design Data for Phenolic Engine Components, Polymer Composite for Automotive Applications: International Congr. Expos., Detroit, Feb 29-March 4, 1988, Society of Automotive Engineers, p. 79–83. 17. J.W. Berg et al., H-SMC Bumper Beam Endured 5 mph Impact, Plastics in Automobiles: Bumper Systems, Interior Trim, Instrument Panels, and Exterior Panels, Society of Automotive Engineers, Detroit, 1989, p. 47–63. 18. J. J. Plomer and T. Traugott, New High Heat Stable, Low Gloss, Automotive Interior Trim Resins Having Excellent Processability, Plastics in Automobiles: Bumper Systems, Interior Trim, Instrument Panels, and Exterior Panels, Society of Automotive Engineers, Detroit, 1989, p. 103–113. 19. N. Pottish, Carbon fiber race car technology hits the streets, High Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, July 2005, p. 52. 20. S. Black, Innovative composite design may replace aluminum chassis, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, February 2006, p. 44. 21. D. Brosius, Carbon fiber goes mainstream automotive, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, April 2003, p. 32. 22. G. Gardiner, Thermoformable composite panels, Part I, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, April 2006, p. 36. 23. D. Brosius, Natural fiber composites slowly take root, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, February 2006, p. 32.

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24. D. Brosius, Detroit auto show a showcase for new composites applications, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, February 2006, p. 20. 25. D. Brosius, NAIA show highlights, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, February 2008, p. 25. 26. P. Malnati, Composites lessen load in fuel-cell demonstrator, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, February 2008, p. 62. 27. Bor Z. Jang, Advanced Polymer Composites, ASM International, Materials Park, 1994. 28. D. Brosius, Composites ousts magnesium in big volume engine, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Feb. 2008, p. 32. 29. Great Britain II, Designed by Alan Gurney, Yachting World, Vol 125 (No. 2772), July 1973, p. 79–80. 30. D. W. Edgell, “Glass Reinforced Plastic Boat Building in the United Kingdom,” Paper 21-D, presented at the 35th Annual Reinforced Plastics/Composites Institute Conference, New Orleans, Society of the Plastics Industry, Feb. 1980. 31. S. Holmes, Resources: Materials: The Lightweight Heavyweight, Insight, Vol 5 (No. 69), Dec. 1981, p. 1902–1907. 32. E. M. Trewin, “Carbon Fibres Performance and Versatility in New Applications,” Paper 23, presented at the Tenth Congress, British Plastics Group, Brighton, Nov. 1976, p. 221–226. 33. P.E. Morgan, E.M. Trewin, and I.P. Watson, Some Aspects of the Manufacture and Use of Carbon Fibre Pultrusion, in Symposium on Fabrication Techniques for Advanced Reinforced Plastics, IPC S&T Press, 1980, p. 69–90. 34. J. Clarke, Partners for the Pond, Yatching World, Vol. 133 (No. 2866), June 1981, p. 70–75. 35. T. Jeffrey, First of the Summer Wine, Yatching World, Vol. 136 (No. 2907), Nov 1984, p. 29. 36. T. Jeffrey, The Colt and the Lion, Yatching World, Vol. 136 (No. 2900), April 1984, p. 98–99. 37. K. L. Pittman, Breaking the Old Moulds, Sail, Jan 1985, p. 76–81. 38. J. Shuttleworth, Spectrum 42 Production Cat, Multihull Int., Vol. 18 (No. 207), April 1985, p. 91–94. 39. E.W. Sponberg, Carbon Fibre Sailboat Hulls: Optimise the Use of an Expensive Material, Marine Technol., Vol. 23 (No. 2), April 1986, p. 165–174. 40. M. Musselman, Europe’s infusion pioneer simplifies process with bottom up approach, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Oct. 2003, p. 34. 41. D. Dawson, In-mold preforming cost-justifies closed molding, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Oct. 2005, p. 34. 42. K. F. Mason, Carbon composites welcomed abroad, High Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Sept. 2004, p. 30. 43. B. Cracknell, J. McWilliam, B. Axford, K.Rose, E. Warden-Owen, and M. Relling, Laminate Sails, Yatching World, Vol. 136 (No. 2009S), April 1984, p. 69–71. 44. K. Rose, Lighter and Stronger, Seahorse, July/Aug. 1984, p. 23, 26. 45. N. Thornton, Yarns From the Trade: Developments in High-Tech Sailcloth, Yachts Yachting, 22 March 1985, p. 26–28. 46. Chain Store Sailmaker, Yachting World, Vol. 137 (No. 2916), Aug. 1985, p. 82–83. 47. J. Daniels, Putting the Sail in Sailboard, Plast. Today, No. 25, Spring 1986, p. 15–18. 48. J. Summerscales, Marine Applications, Engineered Materials Handbook, Vol. 1, Composites, p. 836–844. 49. D. Dawson, New dimensions in sailings, High Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, September 2005, p. 44. 50. G. Bonavent, M. Peniado, Using composite materials in offshore applications, Ocean Ind., Vol. 14 (No. 4), April 1979, p. 372–384. 51. C.P. Sparks, P. Odru, H. Bono, G. Metivaud, Mechanical testing of high-performance composite tubes for TLP production risers. in Proceedings of offshore technology conference, OTC 5797; 1988. p. 467–72. 52. S. Black, Offshore applications: the future is now, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, April 2003, p. 17. 53. O.O. Ochoa, M.M. Salama, Offshore composites: transition barriers to an enabling technology, Composites Science and Technology, 65(2005) p. 2588–2596. 54. D.K. Dawson, Have surfboard, will travel, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, April 2007, p. 54.

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55. C. Red, Carbon fiber sporting goods, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Jan. 2008, p. 34. 56. S. Black, Fundamental fabrication process requires sophisticated design, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Sept. 2004, p. 52. 57. J. Nelson, Aluminum frame build incorporates carbon fiber tubes, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Jan. 2006, p. 44. 58. D.K. Dawson, At the top and still climbing, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, March 2005, p. 46. 59. S. Black, An update on composite tanks for cryogens, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Nov. 2005, p. 22. 60. S. Black, Fiber placement the centerpiece of NCAM, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Jan. 2004, p. 30. 61. B. Berenberg, Tubular composite core stiffens launch vehicle structure, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Sept. 2004, p. 36. 62. B. Berenberg, Towpreg proves cost-competitive for wound pressure vessels, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, July 2003, p. 36. 63. S. Black, Composites alive and well in offshore applications, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, March 2006, p. 34. 64. J. Sloan, In situ manufacturing: Pipe relining in demanding, corrosive environments, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Jan. 2007, p. 34. 65. S. Rush, Dual-laminate pipe comes of age, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, April 2007, p. 36. 66. C. Red, Wind turbine blades: Big and getting bigger, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, June 2008, p. 42. 67. D. K. Dawson, Big blades cut wind energy cost Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Febr. 2003, p. 28. 68. K. F. Mason, Changes in the wind, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, April 2004, p. 26. 69. S. Black, Pultruded composites span infrastructure gaps, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Oct. 2005, p. 28. 70. P. Malnati, New bridge deck bests early FRP systems, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Sept. 2008, p. 54. 71. S. Black, Are hybrid designs the future for composite bridge decks?, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Febr. 2004, p. 20. 72. S. Black, How are composite bridges performing?, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Dec. 2003, p. 16. 73. S. Black, Blast protection for large structures, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, March 2005, p. 56. 74. S. Black, Composites on the frontlines, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Jan. 2005, p. 26. 75. S. Rush, The art of armor development, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, Jan. 2007, p. 22. 76. A. Brown, New options in personal ballistic protection, High-Performance Composites, Design and Manufacturing Solutions for Industry, Ray Publishing, March 2003, p. 23. 77. K.F. Mason, Composites aboard high-speed trains, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Dec. 2004, p. 34. 78. J. Nelson, Passenger safety: flame, smoke and toxicity control, Composites Technology, Engineering and Manufacturing Solutions for Industry, Ray Publishing, Dec. 2005, p. 28. 79. M.M. Davoodi, S.M. Sapuan, D. Ahmad, A. Ali, A. Khalina, M. Jonoobi, Mechanical properties of hybrid kenaf/ glass reinforced epoxy composite for passenger car bumper beam. Mater. Design, Vol. 31, No. 10, Dec. 2010, p. 4927–4932.

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80. R.F. Gibson, A review of recent research on mechanics of multifunctional composite materials and structures. Composite Struct., Vol. 92, No. 12, Nov. 2010, p. 2793–2810. 81. U.A. Khashaba, I.A. El-Sonbaty, A.I. Selmy, A.A. Megahed, Machibability analysis in drilling Woven GFR/epoxy composites: Part II – Effect of drill wear. Composites Part A: Appl. Sci. Manufact., Vol. 41, No. 9, Sept. 2010, p. 1130–1137. 82. M. T. Byrne, Y. K. Gunko, Recent Advances in Research on Carbon Nanotube-Polymer Composites. Adv. Mater., Vol. 22, No. 15, Nov. 2009, p. 1672–1688. 83. R. Stewart, Lightweighting the automotive market. Reinforced Plastics, Vol. 47, No. 2, Aug.-Sept. 2009, p. 902–911. 84. J. Marzbanrad, M. Alijanpour, M.S. Kiasat, Design and analysis of an automotive bumper beam in low-speed frontal crashes. Thin-Walled Struct., Vol. 47, No. 8–9, Aug.-Sept. 2009, p. 902–911. 85. Deo, R.B., Starnes, J.H., Holzwarth, R.C., Low-cost composite materials and structures for aircraft applications. RTO AVT Specialists’ Meeting on “Low Cost Composite Structures”, Loen, Norway, May 2001, http://ftp.rta.nato.int/public//PubFullText/RTO/MP/RTO-MP-069(II)/MP-069(II)-(SM1)-01.pdf. 86. http://www.globalsecurity.org/military/systems/aircraft/f-22-stealth.htm. 87. http://www.vought.com/newsFactGallery/factsheets/programs/787.htm. 88. A. Baker, S. Dutton and D. Kelly, Composite Materials for Aircraft Structures, Second Edition, AIAA Education Series, American Institute of Aeronautics and Astronautics, Reston, VA., p. 1–2. 89. R. Zimmermann and R. Rolfes: Air and Space Europe, Vol. 3, No. 3, May 2001, p. 129–131. 90. I. J. Polmear, G. Pons, Y. Barbaux, H. Octor, C. Sanchez, A. J. Morton, W. E. Borbigge and S. Rogers, Materials Science and Technology 1999, Vol. 15, No 8, p. 861–868. 91. G. Andrei, D. Dima, L. Andrei, Optoelectron. Adv. Mater. 2006, Vol. 8, No 2, p. 726–730. 92. D. K. Hsu, J. J. Peters, F. E. Dong, D. J. Barnard and V. Dayal, SPIE proceedings series. Nondestructive evaluation of ageing of materials and composites. Conference No 3, Newport Beach CA, 1999, Vol. 3585, p. 236–245. 93. T. D’Orazio, C. Guaragnella, M. Leo and P. Spagnolo, NDT and E international 2005, Vol. 38, No 8, p. 665–673. 94. V. K. Varadan and V. V. Varadan, Proceedings of SPIE, the International Society for Optical Engineering, 1999, Vol. 3673, p. 359–368. 95. S. G. Pantelakis and E. A. Baxevani, International conference on fiber reinforced composites No 7, Newcastle upon Tyne, Royaume-Uni, April 1998, p. 208–215. 96. R. Alderliesten, C. Rans and R. Benedictus, The applicability of magnesium based Fibre Metal Laminates in aerospace structures. Composites Science and Technology, 2008, Vol. 68, No 14, p. 2983–2993. 97. H. D. Fair, Symposium on Electromagnetic Launch Technology No 11, Saint-Louis , France, May 2002, IEEE transactions on magnetics, 2003, Vol. 39, No 1, p. 11–17. 98. C. Grant, Automated processes for composite aircraft structure, Industrial Robot 2006, Vol. 33, No 2, pp. 117–121. 99. D. F. Walczyk, J. F. Hosford, and J. M. Papazian, Using Reconfigurable Tooling and Surface Heating for Incremental Forming of Composite Aircraft Parts. J. Manufact. Sci. Engn. 2003, Vol. 125, No 2, p. 333–343. 100. J. Krebs, D. Bhattacharyya, K. Friedrich, Composites. Part A, Applied Science Manufact. 1997, Vol. 28, No 5, p. 481–489. 101. A. A. Baker, P. J. Callus, S. Georgiadis, P. J. Falzon, S. E. Dutton, and K. H. Leong, Composites, Part A, Appl. Sci. Manufact., 2002, Vol. 33, No 5, p. 687–696. 102. http://www.dassaultfalcon.com 103. R. B. Deo, J. H. Starnes and R. C. Holzwarth, Low-cost composite materials and structures for aircraft applications. RTO AVT Specialists “Meeting on Low Cost Composite Structures”, Loen, Norway, 7-11 May 2001; http://ftp.rta.nato.int/public//PubFullText/RTO/MP/RTO-MP-069(II)/MP- 069(II)-(SM1)-01.pdf 104. C. Dry, Proceedings of SPIE, 2008, Vol 6932 (2), pp. 693212.1-693212.13 “Sensors and smart structures technologies for civil, mechanical, and aerospace systems 2008”, San Diego, California, USA 2008. 105. A. Baker, Transactions of the Institution of Engineers, Australia. Mechanical engineering. 1996, vol. 21, No 1-2, pp. 69-87. Symposium on vibration and noise, Newcastle, Australia 1995.

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106. H. S. Choi, K. J. Ahn, J. D. Nam, and H. J. Chun, Composites. Part A, Appl. Sci. Manufact. 2001, Vol. 32, No 5, p. 709–720. 107. H. G. S. J. Thuis, Composite structures 1999, Vol. 47, No 1-4, p. 813-819. Symposium ICCS/10 Intern. Conf. on Composite Structures No 10, Melbourne, Australia 1999. 108. S. Fouquet, M. Rollin, R. Pailler, and X. Bourrat, Nordic Symp. on Tribology, 2008, Vol. 264, No 9-10, p. 850-856. NORDTRIB 2006 No 12, Helsingør, Danemark 2006. 109. P. Linde, and H. De Boer, Modelling of inter-rivet buckling of hybrid composites, Composite structures, 2006, Vol. 73, No 2 pp. 221-228, Intern.Conf. on buckling and postbuckling behavior of composite laminated shell structures, Eilat , Israel 2004. 110. O. V. Startsev, A. S. Krotova and L. T. Startseva, Interlayer shear strength of polymer composite materials during long term climatic ageing. Polym. Degrad. Stab. 1999, Vol. 63, No 2, p. 183–186. 111. J. Jedidi, F. Jacquemin and A. Vautrin, Design of accelerated hygrothermal cycles on polymer matrix composites in the case of a supersonic aircraft. Composite Structures, Vol. 68, No 4, May 2005, p. 429–437. 112. P. D. Mangalgiri, Polymer-matrix Composites for High-temperature Applications. Defence Sci. J., Vol. 55, No. 2, April 2005, p. 175–193. 113. H. S. Choi, Y. H. Jang, Bondline strength evaluation of cocure/precured honeycomb sandwich structures under aircraft hygro and repair environments. Composites Part A: Appl. Sci. Manufact., Vol. 41, No. 9, Sept. 2010, p. 1138–1147. 114. J. R. Majhi, R. Ganguli, Helicopter blade flapping with and wothout small angle assumption in the presence of dynamic stall. Appl. Mathemat. Modelling, Vol. 34, No. 12, Dec. 2010, p. 3726–3740. 115. V.N. Burlayenko, T. Sadowski, Effective elastic properties of foam-filled honeycomb cores of sandwich panels. Composite Struct., Vol. 92, No. 12, Nov. 2010, p. 2890–2900. 116. T. D. Breitzman, E.V. Iarve, B.M. Cook, G.A. Schoeppner, R.P. Lipton, Optimization of a composite scarf repair patch under tensile loading. Composites Part A: Applied Science and Manufacturing, Vol. 40, No. 12, Dec. 2009, p. 1921–1930. 117. U. Beier, F. Wolff-Fabris, F. Fischer, J.K.W. Sandler, V. Altst¨adt, G. H¨ulder, E. Schmachtenberg, H. Spanner, C. Weimer, T. Roser, W. Buchs, Mechanical performance of carbon fibre-reinforced composites based on preforms stiched with innovative low-melting temperature and matrix soluble thermoplastic yarns. Composites Part A: Applied Science and Manufacturing, Vol. 39, Issue 9, Sept. 2008, p. 1572–1581. 118. T. Ogasawara, T. Ishikawa, T. Yamada, R. Yokota, M. Itoh, S. Nogi, Thermal response and ablation characteristics of carbon fiber reinforced composite with novel Silicon Containing Polymer MSP, Journal of composite materials, 2002, vol. pp. 143–157. 119. G. Hidden, L. Boudou, J. Martinez-Vega, S. Remaury, P. Nabara, Development of nanoparticle-polysiloxane composites for spacecraft applications. Polymer engineering and science 2006, vol. 46, pp. 1079–1084. International Symposium on Polymeric Nanocomposites N◦ 3, Boucherville, Quebec, Canada 28/09/2005. 120. E.J. Friebele, C.G. Askins, A.B. Bosse, A.D. Kersey, H.J. Patrick, W.R. Pogue, M.A. Putnam, W.R. Simon, F.A. Tasker, W.S. Vincent, S.T. Vohra, Optical fiber sensors for spacecraft applications, Smart materials and structures 1999, Vol. 8, pp. 813–838. 121. P.B. Joshi, B.L. Upschulte, A.H. Gelb, B.D. Green, D.M. Lester, I.A. Wallace, D.W. Starrett, D.M. Marshall, Autophagous spacecraft composite materials for orbital propulsion. SPIE proceedings series, 2002, Vol. 4698, pp. 171–179. Industrial and commercial applications of Smart structures technologies. Conference, San Diego CA 18/03/2002. 122. W.P. Schonberg, Protecting spacecraft against orbital debris impact damage using composite materials, Composites. Part A, Applied science and manufacturing, 2000, Vol. 31, pp. 869–878. 123. M. Wang, X. Zhao, Z. Shen, S. Ma, Y. Xing, Effects of plerospheres on the atomic oxygen resistance of a phenolic resin composite, Polymer degradation and stability, 2004, vol. 86, pp. 521–528. 124. K. Krishen, Space applications for ionic polymer-metal composite sensors, actuators, and artificial muscles. Acta Astronautica, Vol. 64, No 11-12, June-July 2009, p. 1160–1166. 125. S. Baksi, M. S. Babu, G. Srikant, S. Biswas, Composites – The wonder material. Odern Plastica and Polymers, June 2005.

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126. S. Pulapura, J. Kohn, Trends in the Development of Bioresorbable Polymers for Medical Applications. J. Biomaterials Appl., Vol. 6, No. 3, March 1992, p. 216–250. 127. M. E. Gomes , R. L. Reis, A. M. Cunha, C. A. Blitterswijk and J. D. de Bruijn, Cytocompatibility and response of osteoblastic-like cells to starch-based polymers: effect of several additives and processing conditions. Biomaterials, Vol. 22, No 13, July 2001, p. 1911–1917. 128. S. R. Sershen, S. L. Westcott, N. J. Halas, J. L. West, , Temperature-sensitive polymer-nanoshell composites for photothermally modulated drug delivery. J. Biomedi. Mater. Res., Part A, Vol. 51, No 3, p. 293–298. 129. H. I. Schlaberg, J. S. Duffy, Piezoelectric polymer composite arrays for ultrasonic medical imaging applications, Sensors and actuators. A, Physical, Vol. 44, No 2, Feb. 1994, p. 111–117. 130. D. P. Martin, S. F. Williams, Medical applications of poly-4-hydroxybutyrate: a strong flexible Absorbable Biomaterial. Biochem. Engng. J., Vol. 16, No 2, November 2003, p. 97–105. 131. F. Chen, Z. C. Wang, C. J. Lin, Preparation and characterization of nano-sized hydroxyapatite particles and hydroxyapatite/chitosan nano-composite for use in biomedical materials, Mater. Lett., Vol. 57, No 4, December 2002, p. 858–861. 132. M. Gerard, A. Chaubey, B. D. Malhotra, Application of conducting polymers to biosensors, Bioelectronics, Vol. 17, No 5, May 2002, p. 345–359. 133. P. A. Gunatillake, R. Adhikari, Biodegradable synthetic polymers for tissue engineering, European Cells and Materials, Vol. 5, May 2003, p. 1–16. 134. D. W. Hutmacher, Scaffolds in tissue engineering bone and cartilage, Biomaterials, Vol. 21, No 24, 15, Dec. 2000, p. 2529–2543. 135. J. Jung, B. Kim, Y. Tak, J. O. Park, Undulatory tadpole r Publication Date: 27-31 Oct. 2003, Vol 3, pp. 21332138, IEEE Explore, Intelligent Robots and Systems, October 2003. (IROS 2003). Proceedings 2003 IEEE/RSJ International Conference on. 136. H. Skinner, Composite Technology for Total Hip Arthroplasty. Clinical Orthopaedics and Related Research, Vol 235, October 1988, pp. 224–236. 137. K. Rezwan, Q.Z. Chen, J.J. Blaker, A. R. Boccaccini, Biodegradable and bioactive porous polymer/inorganic composite scaffolds for bone tissue engineering, Vol. 27, No 18, June 2006, p. 3413–3431. 138. B. Adhikari, S. Majumdar, Polymers in sensor applications, Progress in Polymer Science, Vol. 29, No 7, July 2004, p. 699–766. 139. C. T. Laurencin, A. M. A. Ambrosio, M. D. Borden, J. A. Cooper, Jr., Annual Review of Biomedical Engineering, Vol. 1, August 1999, p. 19–46. 140. C. M. Agrawal, R. B. Ray, Biodegradable polymeric scaffolds for musculoskeletal tissue engineering, J. Biomedical Mater. Res., Part A, Vol. 55, No 2, p. 141 – 150. 141. Liac, W.Wang, M. Uo, S. Ohkawa, T. Akasaka, K. Tamura, F. Cui, F. Watari, Biomaterials, Vol 26, No 36, Dec. 2005, p. 7564–7571. 142. M. R. Kim, T. G. Park, Journal of Controlled Release, Vol. 80, No 1-3, April 2002, p. 69–77. 143. M. Yu, H. Shen, Z. D. Dai, Manufacture and Performance of Ionic Polymer-Metal Composites, J. Bionic Engn., Vol 4, No 3, Sept. 2007, p. 143–149. 144. M. Shanipoor, Y. Bar-Cohen, T. Xue, J.O. Simpson, J. Smith, Ionic Polymer-Metal Composites (IPMC) As Biomimetic Sensors and Actuators, Proceedings of SPIE’s 5, Annual International Symposium on Smart Structures and Materials, 1-5 March, 1998, San Diego, CA. Paper No. 33 24–27. 145. M. Shahinpoor, Mass Transfer Induced Hydraulic Actuation in Ionic Polymer-Metal Composites. J. Intelligent Mater. Syst. and Structures, Vol. 13, No. 6, 369–376 (2002). 146. M. Wang, Developing bioactive composite materials for tissue replacement. Biomaterials, Vol. 24, No 13, June 2003, p. 2133–2151. 147. C. Sanchez, B. Juli´an, P. Belleville, M. Popall, Applications of hybrid organic–inorganic nanocomposites. J. Mater. Chem., 2005, Vol. 15, p. 3559–3592. 148. A. P. Marques, R. L. Reis, J. A. Hunt, The biocompatibility of novel starch-based polymers and composites: in vitro studies, Biomaterials, Vol. 23, No 6, March 2002, p. 1471–1478. 149. K. Friedrich, Z. Zhang, A. K. Schlarb, Effects of various fillers on the sliding wear of polymer composites. Composites Science and Technology, Vol. 65, No 15-16, Dec. 2005, pp. 2329–2343.

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150. L. Calandrelli, B. Immirzi, M. Malinconico, M. G. Volpe, A. Oliva, F. Della Ragione, Preparation and characterisation of composites based on biodegradable polymers for “in vivo” application. Polymer, Vol. 41, No 22, Oct. 2000, p. 8027–8033. 151. K. J. L. Burg, S. Porter, J. F. Kellam, Biomaterial developments for bone tissue engineering. Biomaterials, Vol. 21, No 23, December 2000, p. 2347–2359. 152. S. J. Kalita, S. Bose, H. L. Hosick, A. Bandyopadhyay, Development of controlled porosity polymer-ceramic composite scaffolds via fused deposition modeling. Vol 23, No 5, October 2003, pp. 611–620. 153. A. M. Moursi, A. V. Winnard, P. L. Winnard, J. J. Lannutti, R. R. Seghi, Enhanced osteoblast response to a polymethylmethacrylate–hydroxyapatite composite. Biomaterials, Vol. 23, No 1, Jan. 2002, p. 133–144. 154. M. S. A. Bakar, P. Cheang, K. A. Khor, Mechanical properties of injection molded hydroxyapatite- polyetheretherketone biocomposites. Composites Science and Technology, Vol. 63, No 3-4, Febr.- March 2003, p. 421–425. 155. P. M. L´opez-P´erez, R. M. P. da Silva, R. A. Sousa, I. Pashkuleva, R. L. Reis, Plasma-induced polymerization as a tool for surface functionalization of polymer scaffolds for bone tissue engineering : An in vitro study. Acta Biomaterialia, Vol. 6, No. 9, Sept. 2010, p. 3704–3712. 156. M.P. Ginebra, M. Espanol, E.B. Montufar, R.A. Perez, G. Mestres, New processing approaches in calcium phosphate cements and their applications in regenerative medicine. Acta Biomaterialia, Vol. 6, No. 8, Aug. 2010, p. 2863– 2873. 157. M. B. Runge, M. Dadsetan, J. Baltrusaitis, A. M. Knight, T. Ruesink, E. A. Lazcano, L. Lu, A. J. Windebank, M. J. Yaszemski, The development of electrically cionductive polycaprolactone fumarate-polypyrrole composite materials for nerve regeneration. Biomaterials, Vol. 31, No. 23, Aug. 2010, p. 5916–5926. 158. S. Sundelacruz, D. L. Kaplan, Polymeric materials for bone and cartilage repair. Progress in Polymer Science, Vol. 35, No. 4, April 2010, p. 403–440. 159. R. A. Green, N. H. Lovell, G. G. Wallace, L. A. Poole-Warren, Conducting polymers for neural interfaces: Challenges in developing an effective long-term implant. Biomaterials, Vol. 29, No. 24-25, Aug.-Sept. 2008, p. 3393–3399. 160. S. G. Lee, H. C. Park, S. D. Pandita, Y. Yoo, Improvement of IPMC (Ionic Polymer Metal Composites) for a Flapping Actuator, Intern. J. Control, Automation, and Systems, Vol. 4, No 6, Dec. 2006, p. 748–755. 161. T. Christensen, E. Morisbak, H. Hj. Tønnesen, E. M. Bruzell, in vitro photosensitization initiated by camphorquinone nad phanylpropanedione in dental polymeric materials. J. Photochem. Photobiol. B: Biology, Vol. 100, No. 3, Sept. 2010, p. 128–134. 162. N. Moszner, U. Salz, New developments of polymeric dental composites. Progress in Polymer Science, Vol. 26, No 4, May 2001, p. 535–576. 163. J. W. Stansbury, M. Trujillo-Lemon, H. Lu, X. Ding, Y. Lin, J. Ge, Conversion-dependent shrinkage stress and strain in dental resins and composites. Dental Materials, Vol 21, No 1, Jan. 2005, p. 56–67. 164. J. W. Stansbury, Curing Dental Resins and Composites by Photopolymerization. J. Estetic Restorative Dentistry, Vol. 12, No 6, July 2007, p. 300 – 308. 165. K. Fujihara, K. Teo, R. Gopal, P. L. Loh, V. K. Ganesh, S. Ramakrishna, K. W. C. Foong, C. L. Chew, Fibrous composite materials in dentistry and orthopaedics: review and applications, Compos. Sci. Technol., Vol. 64, No 6, May 2004, p. 775–788. 166. M. Atai, M. Nekoomanesh, S. A. Hashemi, S. Amani, Physical and mechanical properties of an experimental dental composite based on a new monomer. Dental Materials, Vol. 20, No 7, Sept. 2004, p. 663–668. 167. J. Janata, M. Josowicz, Conducting polymers in electronic chemical sensors. Nature Materials, Vol. 2, 2003, p. 19–24. 168. S. Black, Electronic applications – Update: Composite PCB. Composites Technology, June 2006, Vol. 12, No. 3, pp. 28–29. 169. Y. Rao, S. Ogitani, P. Kohl, C. P. Wong , Novel polymer-ceramic nanocomposite based on high dielectric constant epoxy formula for embedded capacitor application. J. Appl. Polym. Sci., Vol. 83, No 5, pp 1084–1090. 170. M. S. Dresselhaus, G. Dresselhaus, P. Avouris, Carbon Nanotubes: Synthesis, Structure, Properties and Applications. Topics in Applied Physics 80; Springer: New York, 2001. 171. M. Gerard, A. Chaubey, B. D. Malhotra, Application of conducting polymers to biosensor. Biosensors and Bioelectronics, Vol. 17, No 5, May 2002, p. 345–359.

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172. S. Koul, R. Chandra, S. K. Dhawan, Conducting polyaniline composite for ESD and EMI at 101 GHz Polymer, Vol. 41, No 26, 15 Dec. 2000, p. 9305–9310. 173. O. Savadogo, Emerging membranes for electrochemical systems, Part II. High temperature composite membranes for polymer electrolyte fuel cell (PEFC) applications. J. Power Sources, Vol. 127, No 1-2, March 2004, p. 135–161. 174. F. Liu, B. Yi, D. Xing, J. Yu, H. Zhang, Nafion/PTFE composite membranes for fuel cell applications. J. Membr. Sci., Vol. 212, No 1-2, Febr. 2003, p. 213–223. 175. Q. Xiao, X. Zhou, The study of multiwalled carbon nanotube deposited with conducting polymer for supercapacitor. Electrochimica Acta, Vol. 48, No 5, Jan. 2003, p. 575–580. 176. B. Adhikari, S. Majumdar, Polymers in sensor applications. Progress in Polymer Science, Vol. 29, No 7, July 2004, p. 699–766. 177. P. C. Ramamurthy, A. M. Malshe, W. R. Harrell, R. V. Gregory, K. McGuire, A. M. Rao, Polyaniline/singlewalled carbon nanotube composite electronic devices. Solid-State Electronics, Vol. 48, No 10-11, Oct.-Nov. 2004, p. 2019–2024. 178. R.K. Gupta, R.A. Singh, Fabrication and characteristics of Schottky diode based on composite organic semiconductors. Composites Sci. Technol., Vol. 65, No 3-4, March 2005, p. 677–681. 179. K. Nielsch, F. J. Casta˜no, C. A. Ross, and R. Krishnan, Magnetic properties of template-synthesized cobalt/polymer composite nanotubes. J. Appl. Phys. Vol. 98, No 3, 2005, p. 318. 180. F. Croce, F. S. Fiory, L. Persi, B. Scrosati, A High-Rate, Long-Life, Lithium Nanocomposite Polymer Electrolyte Battery. Electrochem. Solid-State Lett., Vol. 4, No 8, pp. A121–A123 (2001). 181. C. K. Hong, R. P. Wool, Development of a bio-based composite material from soybean oil and keratin fibers. J. Appl. Polym. Sci., Vol. 95, 2005, p. 1524–1538. 182. J. M. J. Fr´echet, Functional polymers: from plastic electronics to polymer-assisted therapeutics. Progress in Polymer Science, Vol. 30, No 8-9, August-September 2005, p. 844–857. 183. R. Nohria, R. K. Khillan, Y. Su, R. Dikshit, Y. Lyoy, K. Varahramyan, Humidity sensor based on ultrathin polyaniline film deposited using layer-by-layer nano-assembly. Sensors and Actuators B: Chemical, Vol. 114, No 1, March 2006, p. 218–222. 184. N. Parvatikar, S. Jain, S. Khasim, M. Revansiddappa, S.V. Bhoraskar, M.V.N. A. Prasad, Electrical and humidity sensing properties of polyaniline/WO composites. Sensors and Actuators B: Chemical, Vol. 114, No 2, April 2006, p. 599–603. 185. M. A. Ryan, A. V. Shevade, H. Zhou, M. L. Homer, Polymer-carbon black composite sensors in an electronic nose for air-quality monitoring. MRS bulletin/Materials Research Society, Vol 29, Oct. 2004, p. 714–719. 186. B. Fugetsu, E. Sano, H. Yu, K. Mori, T. Tanaka, Graphene oxide as dyestuff for the creation of electrically conductive fabrics. Carbon, Vol. 48, No. 12, Oct. 2010, p. 3340–3345. 187. T. H. Nguyen, K. H. Lee, B. T. Lee, Fabrication of Ag nanoparticles dispersed in PVA nanowire mats by microwave irradiation and electro-spinning Mater. Sci. Engn.: C, Vol. 30, No. 7, 30 Aug. 2010, p. 944–950. 188. H. I. Labouta, M. Schneider, Tailor-made biofunctionalized nanoparticles using layer-by-layer. Intern. J. Pharm., Vol. 395, No. 1-2, 16 Aug. 2010, p. 236–242. 189. W. Xu, T. J. Lu, F. Wang, Effects of interfacial properties on the ductility of polymer-supported metal films for flexible electronics. Intern. J. Solids Struct., Vol. 47, No. 14–15, July 2010, p. 1830–1837. 190. Y. S. Kim, Fabrication of carbon black-polymer composite sensors using a position-selective and thickness-controlled electro spray method. Sensors and Actuators B: Chemical, Vol. 147, No. 1, May 2010, p. 137–144.

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24 Waste Management, Recycling and Regeneration of Filled Polymers Jos´e-Marie Lopez Cuesta, Didier Perrin, and Rodolphe Sonnier Ecole des Mines d’Al`es, Al`es, France

24.1 Introduction The end-of-life of polymer materials has become a great challenge, due to environmental concern and to international and national regulations. Among all the possible ways to manage polymer waste, a hierarchy could be established. The most preferred option is the minimization of waste, followed by reuse of materials in the same application, recycling in another application (including recovery of monomers or low-weight molecules), incineration with energy recovery and finally incineration without energy recovery or landfilling [1] (Figure 24.1). But, for a given material, the best option must be tailored. For example, Molgaard shows that recycling of plastics from Danish municipal waste is environmentally sound only if they are separated. If not, incineration with energy recovery is a better environmental option [2]. Life-cycle assessment could help to determine the best waste management route. Composites are often considered as nonrecycling materials whereas national and European directives about environmental impacts of products are more and more stringent. For example, European Directive 2000/53/EC concerning End-of-Life Vehicles (ELV), mandates that 95 wt% of an ELV shall be recoverable by 2015. This directive could slow down the increasing use of composites in the automotive industry. To promote cost-effective recycling solutions, the European Composites Industry Association (EuCIA) and some big European companies founded in 2003 the European Composites Recycling Company (ECRC) [3]. ECRC is developing a Green Label scheme. To obtain the Green Label, fiber-reinforced plastics (FRP) manufacturers must work together to develop recycling routes [4]. Moreover, EuCIA asked the European Parliament to extend the meaning of recycling to the recovery of energy if it is an inseparable part of the recycling process. The goal is to consider the disposal of composites waste by the cement kiln as a recycling option (5).

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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Minimization of waste Reuse of materials Material recycling and feedstock recovery Energy recovery

Least desirable

Figure 24.1

Landfilling – Incineration (without energy recovery)

The waste hierarchy. Reprinted from [1]. Copyright (2006) with permission from Elsevier.

Today, recycling solutions exist for particulate and fiber-reinforced composites. Some routes are only tested at laboratory scale, while others are industrially viable. As for the recycling routes of pristine polymers, one has to distinguish processes devoted to thermoset composites which could not be re-melted and those suitable for thermoplastic composites. In the case of pristine polymers, thermoplastics could be reused as they are after a reprocessing step, while thermosets must be recycled as fillers or degraded to recover chemicals or energy. For multiphase materials, such Sheet Molding Compounds (SMC), another option consists to separate polymer resin, fillers and fibers to recycle them separately. But, before any recycling process, identification and sorting are generally necessary.

24.2 Identification and Sorting The sorting of plastics from a waste source containing many different materials is essential to allow highvalue recycling because a polluting minor phase in a recycled polymer could drastically degrade mechanical properties, even at a percentage lesser than 5 wt%. Various sorting techniques are available for discriminating pure plastics, based on their physical properties (float and sink method, electrostatic or triboelectric separation) or chemical structures (spectroscopic identification). Many of these techniques are not available for filled polymers. For example, polypropylene has a density less than 1, so it is floating in the ‘float and sink’ method used with water. But filled polypropylene with talc could have a density more than 1, according to the filler content, and so the material is sinking. Different spectroscopic techniques have been tested to identify plastics: infrared spectra in the Mid-region (400–4000 cm−1 ), Raman, diffuse reflectance, near infrared (4000–10000 cm−1 ). These methods are based on the comparison between the spectrum of the unknown material and a database. Florestan et al. [6] estimated that Raman is a rapid and highly selective method, allowing the identification of fillers in polymers, as TiO2 and CaCO3 in PP. Stuart Williams et al. [7] report the use of Raman probe gun on a disassembly chain of waste electrical and electronic equipment. Nevertheless, spectroscopic techniques (particularly near infrared spectroscopy) are not very effective to identify dark materials, i.e. polymers containing carbon black, which is one of the most usual fillers, particularly for rubbers. Some studies have concerned the efficiency of electrostatic techniques to identify and sort pure polymers. Hearn and Ballard [8] apply a controlled electrostatic charge to items and measure the rate of charge relaxation which is different according to the type of plastic. Matsushita et al. [9] present another method (tribo-electricity): the plastic items are charged by rubbing against each other and fall through an electric field. The polarities of the items depend on their chemical nature. For example, when PP is rubbed with Polyamide (PA), PP is charged negatively and PA positively. So an electrostatic separation becomes possible. When more than two polymers have to be separated, successive separation steps must be performed. A commercial portative sensor called Tribopen (Figure 24.2), developed by Southampton University in collaboration with

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Figure 24.2

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The portative triboelectric pen called Tribopen. Reprinted from [10]. Copyright (1996) IEEE.

Ford and Wolfson Electrostatics, is based on the same principle [10]. The negative or positive charge generated when the plastic is rubbed with the sensor head allows the identification of the plastic between two types. So this sensor is not available when waste is constituted by numerous unknown composites. Another problem is the identification and sorting of composites containing flame retardants (FR). Even though little data are published about this problem, antagonistic effects of different FR systems could be assumed. Then, it would be more valuable to separate materials based on the same polymer matrix but containing different FR. Up to now, no industrial process has been developed to distinguish FR systems in polymers. Nevertheless, technologies exist to identify some FR systems (as bromine-based FR), as energy dispersive X-ray fluorescence spectrometer. As seen, many difficulties limit the efficiency of the identification and sorting of composites waste. But in many cases, if well-determined waste is collected before blending with other sources of waste, the sorting could be avoided. For example, almost 40% of long carbon fiber, pre-impregnated material (called prepreg) is wasted as offcuts during fabrication [11]. Once the sorting has been achieved, several recycling routes are possible. Incineration is not in our concern because inorganic matter (carbon or glass fiber, CaCO3 filler) of composites is lost in this case while organic matrix is recovered only as energy. Moreover, incineration of composites is often not interesting because of large inorganic fraction. For example, the calorific value of SMC (containing roughly 25% of organic resin, 50% of CaCO3 filler and 25% of glass fiber) is very low: less than 7 MJ/kg [12]. The next section concerns a first approach for composites recycling: separation of matrix, fibers and fillers without the degradation of the components and the separated reuse. Most of the cases deal with the recovery of fibers (glass or carbon fibers) because of the high value of these components.

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24.3 Separation of Components The separation of components could be carried out either by using mechanical separation or more properly by using specific solvents. This last possibility concerns only thermoplastic polymers, but a new kind of thermoset (crosslinkable–decrosslinkable thermosets) could also be recycled through the dissolution. Fibers are obviously the first fillers which could be separated from the polymer matrix and reused but, at a laboratory scale, other fillers like flame-retardant additives were extracted using ionic liquids.

24.3.1

Mechanical Separation

In some cases, a grinding process allows to separate polymer and fibers. For example, end-of-life tires are shredded to release metallic parts and fibers from fabric. Metals are removed by a magnetic system, while dust and fibers are sucked up prior to the grinding of rubber in fine granulates. Metals could be recycled, while very few markets are available for fibers. Fibers could be sold for a tire-derived fuel application. Ground rubber particles contain still rubber and filler (especially carbon black) and could be recycled as fillers into a new matrix (asphalt, for example). Kouparitsas et al. [13] have ground three different thermoset composites (glass fiber/polyester, carbon fiber/epoxy and aramide fiber/epoxy) and collected the fiber-rich fraction after sifting. Recycled fibers were incorporated in a new thermoplastic matrix (ionomer or PP) and mechanical properties were tested. The fiber-rich fraction corresponded to 32.2, 47.1 and 53.2% of the total weight; in the three cases, it contained residual polymer but the authors did not measure accurately its amount. The length of recycled fibers was roughly 6 mm for glass fibers and 8 mm for carbon fibers. No data were given concerning the decrease of length due to the recycling process. Mechanical properties of recycled composites containing glass fibers and aramide fibers were found generally equal or higher than those of the same composite containing virgin fibers, provided that virgin fibers having the same length as recycled fibers are chosen. On the contrary, modulus and tensile strength of composites containing recycled carbon fibers were significantly lower in comparison with composites with virgin fibers. The presence of residual polymer on recycled fibers does not seem a problem, except in the case of carbon fibers, but no explanation was given concerning this result. Mechanical recycling processes are also performed at an industrial scale for glass fiber-reinforced composites [14]. SMC and BMC can be ground and classified into several grades according to the size of particles or fibers. Some fractions could be used as filler in replacement of calcium carbonate in new SMC or BMC. More information concerning this recycling process is presented in the last section.

24.3.2

Dissolution of Resin

For thermoplastic composites, one possibility is the dissolution of the matrix without degrading the polymer [15–19]. The composite is shredded in small pieces (to enhance dissolution) and dispersed in a good solvent at a suitable temperature. After dissolution, the solution is filtered to separate fibers or fillers and polymer. A non-solvent is added to the filtrate: the polymer precipitates and could be recovered. The solvent/non-solvent mixture could be reused after distillation. This scheme is applied commercially for the recycling of PVC from different composites (wires, fabrics. . .). This treatment is known as the Vinyloop process and has been developed by Solvay S.A. After pretreatment (washing and shredding) and dissolution in a solvent of PVC, PVC is separated from the other components by different techniques (centrifugation, decantation). These components are then washed by the solvent to

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remove all PVC residues. PVC is precipitated and additives could be added during this step according to later application. Solvent and water are collected and separated: water is treated and solvent is reused. If recycled fibers are incorporated in a different matrix from the original resin, a perfect washing must be performed to remove any residual polymer. Moreover, coupling agents could be also removed with washing. Then an additional surface treatment is necessary to improve interfacial adhesion between fibers and the new matrix. Conversely, if recycled fibers are incorporated in the same matrix, the presence of residual polymer could be interesting. Papaspyrides et al. [16] have studied the influence of a residual polymer fraction on the surface of recycled fibers. After dissolving the matrix (ionomer or low-density polyethylene), glass fibers are washed with different amounts of solvent according to the number of washings. Ionomer is here a copolymer of ethylene and partially neutralized methacrylic acid. Then, glass fibers are incorporated at 40 wt% in the same virgin matrix. In the first case (ionomer/glass fiber), modulus and tensile strength (at break and at yield) of the new composite are higher than those of the virgin composite when no washing or only one is performed. Remaining polymer on fibers is roughly 1 or 2 wt% in these cases. For 2 or 3 washings, it becomes very low and mechanical properties decrease at the same level as virgin composite although the length of fibers decreases due to recycling. These results are explained by the better dispersion of recycled fibers into the matrix because bundles are able to open during recycling and by a better cooperation between fibers and matrix due to the presence of residual polymer on the fibers surface. Nevertheless, in the second case (LDPE/glass fiber), the remaining polymer does not improve mechanical properties of the recycled composite. In another article, Poulakis et al. [17] removed polypropylene by dissolving it in xylene at 135◦ C and recovered glass fibers. Residual polymer content on fibers is 4% if no additional washing is performed and 1% or even after one or two washings. Recycled fibers were then incorporated into the new polypropylene matrix. Tensile strength and modulus are improved when no washing was performed, which confirms the previous conclusion concerning the beneficial effect of deposited polymer on fibers. For example, modulus is 860 MPa for virgin composite, 930 MPa for recycled composite without washing of fibers, and 900 MPa for one or two washings. On the contrary, impact strength decreases for recycled composite, especially when a high content of polymer is deposited on fibers (Figure 24.3). According to the authors, this result is due to the strong fiber–matrix interfacial bond when residual polymer is present on fiber. Zafeiropoulos et al. [18] have pointed out another phenomenon related to the presence of deposited polymer on fiber surfaces. Depending on processing conditions (especially shear), a transcrystalline layer at the surface of fibers could be initiated by the residual polymer. The size of this layer depends on the amount of residual polymer. The authors suggest that mechanical properties of composites could be strongly influenced by this phenomenon, even if they did not perform any mechanical tests. The dissolution of matrix to recover the minor phase was also carried out by Collier and Baird [19] in the case of a thermotropic liquid crystalline polymer. The composite was constituted of a PP matrix (60 wt%) and a thermotropic liquid crystalline Dupont HX8000 (a blend of terephtalic acid, 4-hydroxybenzoic acid, hydroquinone and its derivates). The first step of the process was a reactive extrusion with dicumyl peroxide to decrease molecular mass of PP by causing chain scissions. In a second step, the blend was extruded into a low viscosity mineral oil. PP was dissolved, but not the HX8000 polymer. Separation was carried out by centrifugation and decantation. HX8000 particles were boiled in kerosene to remove oil. Then the particles were washed with hexane to remove kerosene and dried. 70 wt% of HX8000 could be removed from initial composite and a purity of 97 wt% for the recovered HX8000 fraction was achieved. When recycled HX8000 replaced partially virgin HX8000 into a new composite, mechanical properties were maintained or even increased. An increase was obtained for tensile and flexural modulus and flexural strength when 30 wt% of virgin HX8000 was replaced by recycled HX8000. On the contrary, viscosity of reclaimed HX8000 was significantly lower than virgin HX8000, due to the residual PP.

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MODULUS OF ELASTICITY (MPA)

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0

1 NUMBER OF WASHINGS

2

950 900 850 800 750 1

2

3

4

120 100 80 60 40 1

2

3

4

Figure 24.3 Polypropylene deposited on recycled glass fibers and mechanical properties of recycled glass fiberpolypropylene composites for different grades of glass fibers. 1: Unused glass fibers; 2: recycled glass fibers, no washings; 3: recycled glass fibers, one washing; 4: recycled glass fibers, two washings. Reprinted from [17]. Copyright (1998) with permission from John Wiley & Sons.

The extraction of fillers as components of flame-retardant (FR) systems from polymer matrix is also a challenge because antagonistic effects of different FR systems in recycled materials could probably result in a drastic decrease of fireproofing behavior of the polymer. Moreover, many halogen-based FR are progressively phased out and then the waste management of polymers containing these FR becomes critical. Today, no industrial extraction process has been developed, but some solutions are proposed at a laboratory scale [20, 21].

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Table 24.1 Fractionation of deca-BDE and Sb2O3 between the ionic liquid and ethyl acetate phases in liquid-liquid extraction of flame retardants. Reprinted from [45]. Copyright (2007) with permission from Elsevier. Deca-BDE based on Br content Solid phase Solid precipitated from ionic liquid layer Solid HIPS recovered from ethyl acetate layer

Sb2O3 based on Sb content

Br (%)

Fraction in solid phase

Sb (%)

Fraction in solid phase

52.0

92.7

18.6

99.9

4.1

7.3

0.02

0.1

Lateef et al. [20] have studied the feasibility of the extraction of polybrominated FR and antimony trioxide from electrical and electronic equipment plastics (especially high impact polystyrene) by using ionic liquids. Several ionic liquids were synthesized in a microwave reactor. The best results were obtained when HIPS is finely ground and dispersed at 8% in ethyl acetate at 77◦ C for 10 min. The ionic liquid added was 1-hexylpyridinium bromide. Two layers could be observed. The upper was ethyl acetate and the lower the ionic liquid. After decantation and evaporation of ethyl acetate, methanol was added to ionic liquid to form a precipitate which was identified as a mixture of deca-bromodiphenyl ether (deca-BDE) and antimony trioxide. The results showed that 92.7% of deca-BDE and 99.9% of antimony trioxide were removed from HIPS (Table 24.1). Moreover, no decrease in polymer length was observed, indicating that no degradation occurred. Altwaiq et al. (21) have proposed other processes to extract flame-retardant systems (deca-BDE, tetrabromobisphenol-A (TBBA) and its derivates) from several polymers (ABS, HIPS, PS, ABS-PC). Four procedures have been attempted. The first is a supercritical carbon dioxide extraction. A glass tube containing ground polymer (0.2 g) was placed inside the high pressure autoclave. Pressure was fixed between 10 and 12 MPa using gaseous CO2 . Temperature was maintained at 60◦ C. In the second process (called modified supercritical carbon dioxide extraction), the same procedure was applied but a small content of co-solvent (10 mL) was added to enhance extraction. Different solvents were tested, including toluene, THF, methanol, propan-1-ol. The third process is a solvent extraction. 0.2 g of polymer was dissolved in 20mL of solvent at 60◦ C over 2 h. The same solvents and others were used. The fourth process is a Soxhlet extraction using propan-1-ol or toluene at 190◦ C for 4 h. The extracted solutions were evaporated to obtain solid substances. TBBA carbonate oligomer was efficiently extracted from ABS using toluene in the four cases. Extraction efficiency was higher than 90%. For other cases, brominated FR, solvent and Soxhlet extractions were generally more efficient than the two first processes. Only one FR (hexabromocyclododecane in HIPS) could not be extracted at a high content, regardless of the process. The extraction content was always lower than 10%. Moreover, the authors checked the high degree of purity for extracted FR. A last route presented here concerns the treatment of new types of thermoset polymers which are usually considered as infusible and insoluble. It aims to develop reworkable crosslinked polymer, which means crosslinkable–decrosslinkable systems [22, 23]. In this case, composites based on these new polymers could be remelted as thermoplastic composites. The matrix could be also dissolved to recover separately polymer and fillers/fibers. In these polymers, some bondings could be selectively broken using heat or chemicals. For example, Shin et al. [22] have developed several thermoset systems using di-epoxyde crosslinkers containing sulfonate esters. Heating could lead to the cleavage of these bonds, allowing the reprocessing of the polymer. In one case, the synthesized polymer is totally soluble below 100◦ C and above 180◦ C and partially crosslinked between these two temperatures (and fully crosslinked at 160◦ C).

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24.4 Feedstock Recycling Recycling needs generally depolymerization/degradation for thermoset composites. Chemical recycling is also an option for some thermoplastic composites based on polycondensates such as poly(ethylene terephtalate) (PET), when products of depolymerization are valuable. Several techniques such as hydrolysis [24, 25], alcoholysis/glycolysis [11, 25, 26], acetolysis [24], aminolysis [26], ammonolysis [27], hydrogenolysis [28], have been performed. In these cases, fillers or fibers could be reincorporated in other polymers, while the feedstock could be reused for regenerating the initial polymer or for another goal. Suitable chemical recycling needs not only a precisely controlled depolymerization of the polymer into valuable monomers or oligomers but also a purification method to separate the products obtained. Moreover, for a given polymer, the nature of products obtained by solvolysis depends on the method used. For example, Karayannidis and Achilias [25] reviewed chemical recycling methods for PET. While alkaline hydrolysis of PET leads to terephtalic acid and ethylene glycol, methanolysis degrades PET into dimethyl terephtalate and ethylene glycol. Glycolysis of PET allows the insertion of ethylene glycol in PET chains to give a substrate for unsaturated polyester synthesis [29]. Braun et al. [28] have successfully performed hydrogenolysis on different thermoset polymers, including composites. In a first experience, one gram of ground epoxy resin was put into an autoclave with 5 g of tetraline (as hydrogen donor). Temperature was fixed in the range 320–400◦ C for 5 h. Up to 360◦ C, the soluble fraction was higher than 90 wt%. The main products obtained were bisphenol-A, phenol, p-isopropylphenol, phtalic anhydride, benzoic acid and benzene. Hydrogenolysis of phenolic, melamine-formaldehyde, epoxy, and unsaturated polyester resins crosslinked with styrene and polystyrene crosslinked with divinylbenzene, were also performed. Various fillers were used, including wood flour, chalk, cellulose, glass fiber and kaolin. In one case, the yield of hydrogenolysis was higher than 100% due to the partial degradation of wood flour. In other cases, the yield was generally higher than 90 wt%. The method seems suitable for all thermosets under investigation. The authors have also shown that the degradation could be accelerated and performed at lower temperatures (300◦ C) when amines are added with hydrogen donor. Polyurethane resin could be also degraded using glycolysis [26, 30]. Lee and Kim have performed this method to recycle water-blown rigid polyurethane foam [30]. Glycolysates contain carbamate-type polyols and toxic aromatic amines. Amines could be converted into polyols using butyl glycidylether as reactant. These products are used for the reprocessing of PU foams. The authors have noticed that the replacement of 30 wt% of virgin PU by recycled PU synthesized with glycolysis products did not deteriorate the properties of the foam. Economy and Andreopoulos [24] have recycled thermoset copolyesters using hydrolysis and acetolysis. Hydrolysis would be the preferred method in the case of composites due to the low viscosity of the aqueous solution which allows an easy separation between resin and filler. But, separation of water-soluble and waterinsoluble fractions of degraded resin is difficult. On the contrary, acetolysis takes place in an autoclave at higher temperature than hydrolysis (240◦ C versus 90◦ C). 1.5 ml of glacial acetic acid is added to 10 g of copolyester. Acetolysis leads to oligomeric products which could be directly reused to synthesize a new copolyester resin. Recycled oligomers are applied on glass fabric as coating and cured at 300◦ C during 2 h. Flexural modulus (from dynamic mechanical analysis) is lower for recycled resin than for virgin resin. Nevertheless, recycled oligomers could replace 50 wt% of virgin oligomers without loss of mechanical properties. Solvolysis could be performed under supercritical conditions to enhance diffusivity of the solvent and to improve degradation of polymer. Kamimura et al. [31] used secondary alcohols in supercritical conditions to convert polyamide (nylon 6) into its caprolactam monomers. Degradation was carried out in an autoclave during 1.5 h. Temperature was fixed at 370◦ C. Pressure was measured about 22 MPa, largely above supercritical conditions of the studied alcohols. Yields could reach 97 wt% with a high level of purity.

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Pi˜nero-Hernanz et al. [11, 32] have recently degraded carbon fiber-reinforced composites under supercritical conditions. In a first study, 0.5 g of composite (epoxy resin containing carbon fiber) was placed in an autoclave with water and hydrogen peroxide to enhance partial oxidation [11]. Temperature ranged from 523 to 673 K. Temperature, time, pressure, composite/solvent ratio and H2 O2 concentration were studied parameters. A maximum yield (defined as eliminated resin) of 79.3 wt% was achieved. Temperature was found as the most influential parameter. Adding potassium hydroxide as catalyst allowed the yield to increase up to 95 wt%. Moreover, recovered carbon fibers were almost resin free. Tensile strength values of recovered fibers were found close to those of virgin fibers (3.6–4 GPa versus 4.1 GPa) and significantly higher than those of fibers reclaimed by other recycling processes (microwave or fluidized bed). In another series of experiments, the authors performed chemical recycling on carbon fiber-reinforced composite using alcohols or acetone as solvents under supercritical conditions [32]. The highest yield (98 wt%) was obtained using 1-propanol as solvent-reagent at 350◦ C for 70 min. The use of potassium hydroxide as catalyst allowed yields close to 95 wt% at lower temperatures (250–300◦ C) to be obtained. Recovered fibers were resin free and exhibited similar mechanical properties as virgin fibers. The main drawbacks of these previous methods concern the economic aspects, in particular the high consumption of energy over a long period: treatment is generally performed at high temperatures and pressures (in the case of supercritical processes) for several hours to attain a complete degradation of polymers. Some authors have proposed to use microwave irradiation as an energy source to enhance solvolysis. For example, glycolysis [33] or alcoholysis [34] of poly(ethylene terephtalate) have been performed using microwave. In [34], PET pellets were stirred in a solution containing alcohols in excess and KOH. It was shown that PET could be totally depolymerized using 1-pentanol after only 360 s. Moreover, the authors indicated that the yield and the purity of recycled terephtalic acid were high.

24.5 Thermal Processes This section is devoted to the recycling of components through thermal decomposition leading to the reuse of fillers and feedstock. Two main routes based on heating are available for the treatment of composites: combustion (under air) and pyrolysis (without air). Pickering et al. [35–37] have extensively studied the recycling of composites using a thermo-oxidative process. A fluidized bed process leads to the decomposition of polymer matrix while fibers and fillers are collected and separated for recycling (Figure 24.4). Oxygen allows the minimization of char formation. A fluidized bed (1 m/s) is controlled to elutriate released fibers and fillers but not non-degraded material. Organic contaminants are volatilized with polymer but inorganic solids (e.g. metal) sink in the fluidized bed. A rotating sieve separator is added to separate long fibers from short fibers and particulate fillers. Temperature is fixed at 450◦ C but a second combustion chamber (at 1000◦ C) leads to the complete combustion of organic volatiles producing a clean flue gas and allowing the recovery of energy. Therefore, in such process, polymer resin is not recycled as feedstock but as energy. Conversely, fibers could be recovered and reused in new composites. Different components of composites were recycled using fluidized bed process and recovered glass fibers have been studied. The fiber fraction from SMC was contaminated by fillers and remaining polymer. The fiber purity is only 66 wt% but contaminant content fell to 10 wt%. The fiber yield was measured at 44 wt% only due to the large content of fines produced during previous size reduction stages. The mean fiber length was approximately 5 mm. Mechanical properties of recovered fibers have been tested. Tensile strength was reduced by 50%, comparing virgin fibers but Young’s modulus was not affected by recycling process. Recovered fibers have been incorporated into a typical Dough Molding Compounds (DMC) formulation for compression molding application. When the fiber fraction contains less than 50% of recovered fibers, tensile, flexural and impact properties are not strongly affected. Above 50% of substitution, all properties fall

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FRP Feed

1200mm Fillers

Fibre collection bin

bed support and air distributor plate

Bed

Pre - heaters Air Inlet

Figure 24.4 Diagram of the fluidized bed processing rig. Reprinted from [35]. Copyright (1998) with permission from Elsevier.

significantly, except Young’s modulus which was not modified (Figure 24.5). The authors considered that this process was economically efficient if 9000 t of scrap could be recycled per year. But improvement of fiber yield could decrease this threshold. Other studies have been carried out to maximize the mechanical properties of reclaimed glass fibers [35, 36]. SMC has been treated using fluidized bed process and glass fibers have been recovered. Temperatures ranged from 450 to 650◦ C and fluidizing velocity ranged from 1.3 to 1.7 m/s. Fibers have been washed to remove contaminants (60% by weight). Ultimate tensile strength of fibers decreased when process temperature

150 Flexural strength (MPa)

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50

0

0

20 40 60 80 Non virgin reinforcement (%)

100

Figure 24.5 Effect on DMC flexural strength of replacing virgin fibers with reclaimed fibers (black point: recovered fibers silane coated). Reprinted from [45]. Copyright (1998) with permission from Elsevier.

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Table 24.2 Product yields for catalytic conversion of automotive scrap. Reprinted from [20]. Copyright (2008) with permission from John Wiley & Sons. Scrap material SMC (sheet molding compound) ASR (auto shredder residue) R-PP (reinforced polypropylene) Mixed body panels

Solid residue (%)

Liquid product (%)

Gaseous product (%)

82 39 37 31

2 35 18 46

16 26 45 23

increased: approximately 1.3 GPa for 450◦ C, 0.7 GPa for 550◦ C and 0.2 GPa for 650◦ C (against more than 2 GPa for untreated fiber). On the contrary, fluidizing velocity had no significant impact in the range tested. Pyrolysis of composites has been studied by several research groups. Generally, three fractions are obtained: gaseous, liquid products and solid residue. Gaseous products can be reused as energy, for example to selfsustaining process. Liquid fraction can be recovered as feedstock or fuel. Nevertheless, this is often a mixture of numerous chemicals and distillation (or another method) is generally needed to separate products into valuable products. Solid residue is constituted of inorganic fibers or fillers and char. High content of char is a main drawback for pyrolysis process. The content and the composition of each fraction depend on the treated material and the pyrolysis conditions (time, temperature, catalyst. . .). Allred and Busselle [38] have studied pyrolysis process of different automotive scraps under a vacuum between 0.06 and 0.07 MPa. Time process has been fixed at 30 min and various temperatures have been tested. A temperature of 300◦ C was considered as an optimum. As polymer depolymerizes, hydrocarbon conversion products are removed as a gas and condensed in condensers. Non-condensable gases are incinerated. Product yields differ strongly from the initial scraps. For example, SMC gives 82% of solid residue (due to high level of fibers/fillers), 2% of liquid product and 16% of gaseous product, while auto shredder residues (ASR) give 39%, 35% and 26% respectively (Table 24.2). Moreover, liquid conversion products are firstly styrene (57.8%) for SMC, while liquid fraction from ASR consists of many chemicals with a percentage lower than 20%. The authors report also that reclaimed glass fibers have lost 50% of strength. Cunliffe et al. [39, 40] have recycled various fiber-reinforced polymeric waste using pyrolysis. In a first study, the authors have determined the composition and the characteristics of different products obtained from SMC pyrolysis [40]. Gaseous products (6 wt% of products) were mainly CO and CO2 . Oil product (40 wt%) has fuel properties (low viscosity, no high content in contaminants) and could be reused as fuel oil even if gross calorific value (GCV = 33.6 MJ/kg) is lower than conventional hydrocarbon liquid fuels, due to high content in oxygen. But oil could be also considered as a feedstock because it contains 26 wt% of styrene. A solid condensable wax (15 wt%) containing mainly phtalic anhydride could be recovered as feedstock too. Recycled glass fibers have been cleaned from char and have been incorporated into a new composite (DMC) to partially replace virgin fibers. The ratio between recycled and virgin fibers is 0.25. Mechanical properties decrease slightly in comparison with a control sample (100% of virgin fibers) but values remain in standard deviations. For example, flexural modulus and Charpy impact strength of control DMC are 9.0 GPa and 15 kJ/m2 respectively against 7.3 and 11 for DMC containing recycled glass fibers. In a more extensive work [39], the authors confirmed that product yields depend greatly on composite composition and process temperature. De Marco et al. [12, 41, 42] have also tested the efficiency of pyrolysis for the recycling of SMC (24.9% of orthophtalic polyester resin containing 25% of glass fibers and 46.7% of calcium carbonate fillers). The authors observed that pyrolysis is not complete when temperature is fixed at 300◦ C. On the contrary,

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no significant difference in pyrolysis products was observed in the range 400–600◦ C. Moreover, at high temperature (600–700◦ C) calcium carbonate fillers were degraded into calcium oxide CaO. Therefore, the most suitable temperatures have been determined between 400 and 500◦ C. According to the authors, gaseous product could be reused to provide energy to the process. Oil fraction has a high GCV (36.8 MJ/kg) and is non-polluting liquid fuel. 40 wt% could be reused as petrol and 60 wt% could be mixed with fuel oils. Solid residue recovered from pyrolysis at 500◦ C contains 65% of calcium carbonate and 35% of glass fibers. The authors noticed that recycled glass fibers are not as good as virgin fibers but they considered that a small content (6 wt%) of solid residue could be incorporated into a new composite BMC (24.9% of polyester resin, 50% of calcium carbonate, 16% of virgin glass fibers and 6% of solid residue) with acceptable mechanical properties. Kaminsky et al. [43, 44] have also studied the pyrolysis of polymers, but in the presence of other kinds of fillers. The recovery of methyl methacrylate (MMA) from PMMA was strongly disturbed in the presence of aluminium tri-hydroxide (ATH), a flame-retardant additive [43]. The blends contained approximatively 66 wt% of ATH and 33 wt% of PMMA. Water released by the decomposition of ATH led to a partial saponification of MMA producing methacrylic acid and methanol. The yield of MMA could be reduced from 90% for pure PMMA to 53% for PMMA-ATH at 450◦ C. The optimization of pyrolysis conditions increased the yield up to 80%. However, due to the large content of ATH, the recovery of MMA was low compared to the material mass. In another study, Kaminsky et al. [44] recycled rubbers (in particular natural rubber containing carbon black from truck tires) by pyrolysis in a fluidized bed. The carbon black is the more valuable product in these materials. The authors observed that the carbon black produced in the range of 700–790◦ C closes the pores of the original carbon black and makes it less valuable. Therefore, they concluded that it is better to perform the pyrolysis at 600◦ C: the carbon black content is lower (30 wt% against 40 wt%) but with a higher quality. Nevertheless, these processes have several main drawbacks. First, as seen, some valuable fillers as glass fibers are very sensitive to heat treatment and lose a significant amount of strength (50%). Secondly, the polymer resin is not recovered as materials but only as energy (in many cases). Thirdly, the volatiles released by heating are often toxic or non-environmentally friendly.

24.6 Mechanical Recycling of Filled Thermoplastics The mechanical recycling of thermoplastic polymers consists, after shredding and grinding of end-life thermoplastic parts, in polymer processing operations in which the polymer is melted again in order to obtain new granulates or end products obtained for example by injection molding procedures. Nevertheless, the nature of polymer materials often made of polymer blends (minor components are often present even after use of effective separation processes), containing fibers and fillers strongly influences the behavior of material during reprocessing and its final properties. Degradation processes can occur during all the reprocessing steps. Moreover, the complex composition of the initial material may need the use of additives and compatibilizer agents to control the final morphology and properties.

24.6.1

Degradation During Reprocessing of Filled Thermoplastics and Influence of Interfacial Agents

Mechanical recycling can cause degradation of the polymer matrix and also in some cases a degradation of additives, fibers and fillers present in the polymer, as seen in Figure 24.6, showing the breaking of glass fibers after successive reprocessing of reinforced polypropylene [45]. At first, size reduction operations for

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0.8

Fibres length (mm)

0.7 0.6 0.5 0.4 0.3 0.2 0.1 0

1

2 3 4 Injection cycles

5

6

7

Figure 24.6 Evolution of glass fiber length with the number of recycling. Reprinted from [45]. Copyright (2007) with permission from Elsevier.

end-of-life plastics lead to polymer flakes, allowing the incorporation in processing devices such as extruders or injection molding devices. The effect of these operations on polymer degradation is generally neglected in comparison with the effect of thermomechanical and thermochemical phenomena occurring in the processing devices. According to the nature of polymer and the reprocessing conditions, chains scissions, crosslinking or elimination reactions as well as oxidation and hydrolysis can lead to the degradation of the polymer structure [46]. Some of the reactions can also occur in a virgin polymer if inadequate processing parameters, such as temperature, are selected, for example the elimination of side groups or atoms (HCl in PVC, acetic acid in Ethylene Vinyl Acetate). Nevertheless, some phenomena can be caused by the presence of fillers and impurities, as well as by the lack or the alteration of stabilizers. Several thermoplastic polymers, such as polyamides and polyesters, are very sensitive to hydrolysis during processing and particular attention has to be paid to the dehydroxylation during processing when fillers present in the material to be recycled contain hydroxyl groups such as oxides, silicates or, above all, hydroxides. Hydrolysis reactions also ascribed to moisture present in impurities and fillers can lead to a reduction of melt viscosity, leading to a large decrease of mechanical properties [47]. The main strategy for maintaining the viscosity during reprocessing is an intensive drying or a reprocessing with degassing vacuum. Moreover, chain extenders can be used [48, 49]. These compounds act as polyfunctional compounds able to react, for example, with the carboxylic end groups of PET, leading to a molecular weight increase of the polymer. They react with the PET in the melt, at processing temperatures in a twin screw extruder. The lack of stabilization is often a source of degradation for the recycled thermoplastics. It can be due to a low amount of stabilizers, alteration of stabilizers, or to their possible adsorption on fillers. A too low level of heat stabilizers or anti-oxidants can be explained by evaporation, exsudation, or migration processes. In some cases, the initial amount of stabilizers introduced is not enough to ensure a reprocessing [50]. Moreover, some additives and fillers present in the materials to be recycled can interfere with the mechanisms involving the stabilizers. The presence of halogenated flame retardants in some recycled polymers could entail a release of brominated radicals during reprocessing which could interact with radical scavengers playing the role of antioxidants. The presence of fillers in the waste polymers can act diversely on the stabilization during

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and after reprocessing. Carbon black used as pigment and heat stabilizer is often present in waste polymers. Despite its negative aspect on the color of recycled materials issued from blended sources, its presence can confer a better resistance to degradation processes, due to its ability to scavenge radicals and play the role of hydrogen supplier. Other kinds of fillers could influence stabilization processes. Augier et al. [51] have shown that wood fiber fillers could accelerate PVC degradation leading to dehydrochlorination, formation of unsaturation and then crosslinking. As a consequence, the polymer chain length increases and the matrix shows better mechanical properties, which enhances the entire mechanical properties of the composite. Conversely, various fillers contain metallic ions as impurities or elements of the crystallographic structure (case of lamellar silicates) and some of these metallic ions, mainly transition metals can catalyze the decomposition of hydroperoxides formed by polymer oxidation, leading to an acceleration of the oxidation of the polymer by promoting chain reactions [52]. Consequently, metal deactivators should be incorporated as additives when the reprocessing of the waste polymer is carried out. The influence of impurities present in fillers on functional properties of filled thermoplastics would concern particularly fillers having a high specific surface area such as submicronic particles. The development of polymer nanocomposites for industrial applications is strongly dependent on the long-term stability of nanometric morphologies regarding thermal and photochemical ageing. Several studies have shown that the incorporation of organo-modified montmorillonite in polypropylene with exfoliated nanocomposite morphologies, entailed adverse effects upon long-term UV stability. The presence of iron impurities in the clay particles was proposed to explain its effects [53]. The use of metal deactivators was also applied successfully. The role of filler surface treatments about polymer degradation during reprocessing is not well known. Nevertheless, it can be expected that reactive surface treatments performed on the functional groups of the filler surface could prevent the negative effect of these groups on the polymer stability during reprocessing. In a recent study, Swoboda et al. [54] have incorporated submicronic kaolin at a percentage of 5 wt% in a recycled PET. It has been observed that a surface treatment using triphenylphosphite grafted on filler hydroxyl groups could significantly improve the tensile modulus of composite in comparison with untreated kaolinite, without apparent modification of filler dispersion. It was concluded that triphenylphosphite behaved as a chain extender; nevertheless, it can be also suggested that the surface treatment could limit the risk of hydrolysis of PET due to dehydroxylation. Surface treatments of fillers, and particularly fibers, as well as the use of coupling agents seem also to act positively on their mechanical degradation during reprocessing. A better adhesion between high aspect ratio fillers and the polymer could limit its breakage during the reprocessing steps. According to Xanthos [55], it is possible, through the addition of glass fibers as reinforcements and in the presence of suitable adhesion promoters such as maleic anhydride grafted polypropylene (PPgMA) or styrene-butadiene-styrene (SEBSgMA) at the polymer/fiber interface, to obtain composites from mixed plastics with enhanced and reproducible properties. The use of interfacial agents is particularly important for the recycling of polymers containing natural fibers such as flax, sisal or hemp. In many cases, PPgMA has also been used to limit their breakage during reprocessing. Bourmaud and Baley [45] have determined the recycling behavior of PP/vegetal fiber composites. Different composites using hemp and sisal were characterized. All results were compared with PP-g-MA/hemp composites and PP/glass fiber composites. The results proved that mechanical properties are well conserved with the reprocessing of PP/vegetal fiber composites but that there was poor adhesion between the fibers and PP without any treatment. A significant decrease of fiber length during reprocessing induced by the injection and mechanical grinding processes was noticed. The use of PPgMA induced an improvement of the adhesion between fibers and PP but this effect disappeared after seven injection molding cycles. The degradation of the interfacial agent has been suggested. Rheological experiments evidenced an important decrease of composite viscosity with the number of cycles, which was ascribed to either chain scissions induced by reprocessing and grinding or fiber length decrease. Moreover, thermal analysis showed

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that vegetal fibers had a nucleating ability of vegetal fibers for the crystallization of PP, sisal or hemp fibers accelerating the crystallization process. Among the other available coupling agents allowing to improve the natural fiber/polymer interface, poly[methylene (polyphenyl isocyanate)] (PMPPIC) has been used by Maldas and Kokta [56] to ensure the feasibility of recycling for polystyrene-hardwood aspen fiber/polystyrene composites. The mechanical properties and dimensional stability of composites were investigated under extreme conditions (exposure to boiling water and at room temperature as well as exposure to +105◦ C and –20◦ C). The influence of the coupling agent at 3 %wt, and various treatments, such as fiber coated with 10% polymer +8% PMPPIC and grafted with polystyrene, on the properties of the composites have also been studied. Compared with the original composites, the mechanical properties and dimensional stability of the recycled composites did not change significantly even after exposure to extreme conditions. Moreover, the treated composites offered improved properties compared with nontreated and original polymer under all experimental conditions. 24.6.2

Properties of Recycled Filled or Reinforced Thermoplastics

Several kinds of recycled filled or reinforced thermoplastics can be distinguished. In a first category, the material is still a particulate composite before recycling, and this section will focus on their possible modification after regeneration and even after successive reprocessing. The other categories correspond to the incorporation of fillers to improve the properties (above all mechanical ones) after recycling. On the one hand, these fillers can be usual fillers, able to produce generally a reinforcement effect; on the other hand, these fillers can also be recycled fillers from diverse origins: waste minerals, ground composites or polymers, agricultural waste. 24.6.2.1

Use of Various Fillers to Upgrade the Properties of Recycled Thermoplastics

Among the recycled fillers introduced in polymers to be regenerated, ground fillers rubber from tires (GTR) have been investigated in many recycled thermoplastic polymers, particularly polyolefins. Nevertheless, various authors [57–59] have shown that GTR were not able to confer either a reinforcement of polymers or an elastomeric behavior. Several methods have been carried out to modify the GTR/polymer to improve the composite properties, namely grafting [60], reactions using peroxides, irradiation of GTR and host polymers [61, 62]. Many crosslinked polymers have been ground and reincorporated in thermoplastic or thermoset polymers as fillers. The reuse of ground polyester composites is developed in the section. The case of polyurethane is also significant and many materials produced from it end up as waste. Cryogenic micronization of PU foams can lead to particles smaller than 100 μm. These particles are dispersed in polyols and this mixture is reacted with isocyanates to produce new foams with an equivalent density, provided that the amount of PU filler is not higher than 5 wt% [63, 64]. Micronized PU has also been incorporated as filler in vulcanized rubber and compared with carbon black by Sombasonpop and Sims [65]. Some minerals powders, by-products of industrial processes can play a role of alternative materials to replace usual fillers in recycled and even in non-recycled materials. For example, fly ash [64] and silica fumes [66] have been used in polypropylene. Fly ash and silica fumes are respectively by-products of thermal power stations and ferrosilicon or silicon production plants. In the work of Lafaurie et al. [66], the use of surface treatments with amines allowed to improve the dispersion of the submicronic particles of silica fumes, leading to an improvement of tensile strength in comparison with pristine PP. The use of reclaimed natural fibers (wood pulp, sisal, hemp, flax. . .) is increasing due to their low cost, availability as renewable resources and biodegradability. Some natural fibers seem able to confer reinforcements similar to those conferred by glass fibers [67]. Nevertheless, comparisons made by Bourmaud

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Table 24.3 Mechanical properties of fiber-reinforced composites. Reprinted from [38]. Copyright (2000) with permission from SAGE Publications. Materials characteristics from producers Material PP/hemp 70/30 PP/sisal 70/30 PP-MA/hemp 70/30 PP PP/glass 70/30 Material PP/hemp 70/30 PP/sisal 70/30 PP-MA/hemp 70/30 PP PP/glass 70/30

Tensile modulus (MPa)

Elongation at break (%)

Tensile strength at yield (MPa)

4100 3150 3100 1600 6900

2.0 3.0 4.5 148.0 2.0

32.0 26.6 32.0 26.0 73.5

Density at 25◦ C (g/cm3 )

Glass transition (◦ C)

Flow rate (2.16 kg and 230◦ C) (g/10 min)

0.94 0.99 0.94 0.92 1.16

−9.5 −9.5 −9.5 −6.0 −10.0

1 3 1 18 12

and Baley [45] about mechanical properties of recycled PP reinforced by hemp of sisal and glass fibers at 30 %wt show significant higher modulus and tensile strength for glass fibers (Table 24.3). Reclaimed glass or carbon fibers as industrial scrap of the composites industry can interestingly play a role of reinforcement for different polymers according to their type of surface treatment. Drozdov et al. [68] have reported on virgin, recycled and a mixture of virgin and recycled polycarbonates reinforced with various amounts of short glass fibers in tensile tests with a constant strain rate and in oscillatory torsion tests at room temperature. The study has focused on the effects of filler content and recycling of the host matrix on the mechanical properties of composites. The Young modulus, and the shear modulus, of a polymer composite linearly increased with the filler content and only slight differences were noticed in relation with the matrix composition. It was concluded that the recycling of polycarbonate weakly affects the viscoplastic behavior of polymer composites filled with short glass fibers, but noticeably influences their viscoelastic response. In some cases, polymers to be recycled are complex mixtures of various polymer wastes or contain a significant fraction of a minor polymer contaminant. Due to the absence of miscibility for the most part of polymer blends, strategies of compatibilization have been developed to improve interfacial adhesion. In addition, the use of fillers can be expected to upgrade the mechanical properties of the blends. Xanthos [55] has studied multiphase, multicomponent polymer mixtures equivalent to those found in commingled waste streams, such as those obtained from reclamation/recycling operations of postconsumer containers. The effects of variations in the composition of matrices containing HDPE as the major phase (80 %wt), PS, LDPE, PET or PVC as the minor phases, on the properties of composites containing varying amounts of glass fiber (10–20 %wt) and different adhesion promoters have been studied. The use of both these last interfacial agents and glass fibers allowed strong improvements of tensile strength and modulus to be produced without dramatic losses of elongation at the maximal resistance and the impact resistance of unnotched samples was maintained in comparison with the reference polymer blends. Moreover, relatively little dependence

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of mechanical properties on matrix composition for reinforced thermoplastic blends with additives was noticed On the whole, the use of fillers to improve the microstructure and to maximize the mechanical properties of recycled polymer blends requires the filler/polymer interface(s) to be considered. In some cases, compatibilizing agents can act both at polymer/polymer and filler/polymer interfaces. Grafted polyolefins or olefinic elastomers with maleic anhydride can compatibilize polymer blends due to the polar character of this functional group able to interact with polar polymers such as polyamides, for example, and the apolar character of the chain which ensures a good interaction with polyolefins. Moreover, the maleic anhydride group is able to create ionic links with some kinds of filler, mainly carbonates. Sahnoune et al. [69] have investigated the behavior of HDPE/PS blends and studied the combined influence of interfacial agents and calcium carbonate on microstructure and mechanical properties of these blends. The incorporation of CaCO3 alone significantly enhanced the stiffness but lowers the impact resistance. SEBS (styrene ethylene butylene styrene) and SEBSgMA (grafted with maleic anhydride) elastomer copolymers were incorporated in order to compensate for the embrittlement caused by the CaCO3 filler. Depending on their chemical structure, either grafted with a reactive function or ungrafted, the elastomers acted differently at the interfaces of the HDPE/PS/CaCO3 system. It was observed that SEBS acts exclusively at the HDPE-PS interface whereas SEBSgMA acts at both the HDPE-PS and the HDPE-CaCO3 interface. On the one hand, the SEBSgMA elastomer lowered the stiffening effect caused by CaCO3 and provided an insufficient increase in impact properties. On the other hand, SEBS, which concentrates its action at the HDPE/PS interface, retained much of the stiffening effect of CaCO3 and provided a greater improvement in impact properties than SEBSgMA. In the recycled HDPE/PS (75/25) blend, the incorporation of 20 wt% CaCO3 and 4 wt% SEBS led to an increase in both impact strength (from 39 to 70 kJ/m2 ) and in stiffness (from 1335 to 1560 MPa). 24.6.2.2

Performances of Recycled Thermoplastic Composites

The main research works presented in the literature concern recycled thermoplastic containing glass or natural fibers. The properties generally investigated are the mechanical ones. In many cases, the experiments carried out have been performed using virgin polymers and not reclaimed materials. The methodology of all these studies was generally based on multiple reprocessing of the filled polymers using extrusion and/or injection molding procedures, as it was previously developed for the study of the recycling of unfilled polymers. The performances of recycled glass reinforced polyamides were studied by Eriksson and al. [70]. The authors have focused on the influences of ethylene propylene diene monomer (EPDM) impurities on the durability and the reliability of the mechanical properties of recycled glass fiber-reinforced polyamide 66. A critical size-concentration zone could be determined based on changes in tensile properties of samples containing untreated glass beads of different size and concentrations, which were used as simulated impurities in reprocessed material. Characterization of samples based on recyclate from post-use radiator endcaps, which contained residues of EPDM rubber, showed that the tensile strength was almost unaffected by rubber particles up to an equivalent diameter d, around 700 μm. However, a significant strength reduction commenced for rubber particles larger than 1200–1300 pm, and the strength showed a correlation with d−l/2 . Another study of the same authors [71] have investigated the influence of successive recycling processes on fiber shortening and on the short-term performance of recycled samples compared to that of virgin samples. The results indicated that fiber shortening had a strong influence on strength. Applying a modified Kelly-Tyson model to the fiber length distribution gave excellent agreement with measured strength. There was no need to vary interface or matrix properties in the theoretical analysis. The effect of reprocessing on these factors did not appear to influence strength within the bounds of the model. The decrease in strength during a continuous recycling process was small at a 30 wt% regrind level. Indeed, below 50 wt% regrind, the strength remained

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Figure 24.7 Impact strength as a function of reciprocal number average fiber length. Reprinted from [70]. Copyright (1996) with permission from John Wiley & Sons.

within design limits. The impact strength of dry unnotched samples indicated that the resistance was related to the reciprocal fiber length (Figure 24.7). The effect of mechanical recycling upon tensile strength of an injection molded polyamide 6-6 reinforced with 35% by weight of glass fibers has been also experimentally investigated by Bernasconi et al. [72]. The main effect of recycling was ascribed to fiber breakage with consequent lowering of fibers’ contribution to composite strength. The results from the experimental tests have been compared with predictions obtained by applying a micro-mechanical model, which allowed taking into account the fiber length distribution and the properties of the phases of the composite. Licea-Claverie et al. [73] have studied the recycling of nylon 6-6 based composites with mixed glass fiber and carbon fiber reinforcements, including nylon composite scrap in their formulation Several formulations were characterized by stress–strain measurements, impact testing, and thermal analysis. No dependence on mechanical properties due to increasing amounts of scrap in the composites was found up to 10 wt%. The recycled composites generally showed lower mechanical properties as compared with the virgin composite because of a poor matrix–fiber adhesion. The recycling of other glass fiber-reinforced engineering polymers such as polyphenylene sulfide (PPS) and polycarbonate (PC) has also been studied. Papanicolaou et al. [74] have investigated the effect of hygrothermal aging on the impact strength of virgin and recycled commercial PPS matrix composites reinforced with 40% by weight of short glass fibers after water absorption at different pH and temperature environments. The results showed a great effect of the parameters studied, such as water temperature, alkaline or acidic nature of the environment, and the type of coupling agent used (aminosilane or zirconate) on the impact behavior of the PPS-glass fiber composites. In the report of Chu and Sullivan [75], the recyclability of a fiber-reinforced cyclic bisphenol A polycarbonate has been studied. It has been found that recycled composite using injection and extrusion compression molding yield recycled composites with good tensile properties, though the impact strengths were relatively low. This was ascribed to high fiber orientation and fiber bundle dispersion. On the other hand, compression molded samples, which show random fiber orientation and low fiber bundles dispersion, have relatively low tensile properties, but excellent impact strength. Results have been discussed in relation with the microstructure, including resin molecular weight and fiber length and orientation.

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Natural fibers are frequently used in thermoplastic polymers due to their interesting cost, regarding their reinforcement potential. As for glass fibers, it can be expected that the capacity to recycle this class of composites would particularly depend on the effect of fiber breakage during recycling upon mechanical properties. The recycling of wood fiber-reinforced PVC composite was investigated and compared by Augier et al. [76]. Twenty extrusion-milling cycles were performed and the mechanical and thermal properties evaluated. The mechanical processing broke down the fibers especially during the first cycle, but the wood fibers aspect ratio remained constant. Up to five cycles, the composite properties remained stable. But after 10 cycles, the flexural strength increased, whereas the other mechanical properties remained almost constant. The authors have shown that increase in mechanical properties was only dependent on the influence of the fibers, which accelerated the PVC degradation, characterized by dehydrochlorination followed by crosslinking reactions. The aim of the work of Arbelaiz et al. [77] was to compare the effect of modification routes of short flax fiber bundle/polypropylene (PP) composites on mechanical properties. Modifications were carried out on fiber surface and also modifying PP matrix using several amounts of maleic anhydride PP copolymer (PPgMA) as compatibilizer. After passing MAPP-modified flax fiber bundle/PP composites several times through an injection molding machine, mechanical properties only slightly changed (Figure 24.8). Wielage et al. [78] have studied specific properties of flax and hempfiber- reinforced PP by means of dynamic–mechanical analysis (DMA). The recycling behaviour of natural fiber-reinforced PP showed that multiple processing had only an insignificant influence on the fiber lengths and the mechanical properties. In addition, it was possible to very quickly draw conclusions about the quality of the composite material, such as fiber–matrix adhesion and damping behavior. Fractographic evaluations in the scanning electron microscope (SEM) have confirmed the quantitative characterization obtained from DMA. Due to their biocompostability, natural fiber-reinforced biopolymers, also known as biocomposites, become an attractive alternative to glass fiber-reinforced usual polymers. The ability of flax/PLLA (poly(L-lactide)) biocomposite (20% and 30% fibers by weight) to be recycled has been investigated by Le Duigou et al. [79]. Mechanical properties were evaluated initially, and shown to be similar to those of glass/PP and superior to hemp/PP and sisal/PP composites. After repeated injection cycles, tensile properties were shown to be conserved until the third cycle (Figure 24.9). Experimental techniques revealed a lower molecular weight, lowering of glass transition temperature, reduction of fiber length, and separation of fiber bundles with

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Figure 24.8 Tensile behavior of 5 %wt MAPP modified flax fiber bundle/PP composites as impact strength as a function of fiber loading and number of passings: (a) strength, (b) modulus. Reprinted from [77]. Copyright (2005) with permission from Elsevier.

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Figure 24.9 Evolution of tensile strength at yield of Flax/PLLA as a function of injection cycles. Reprinted from [79]. Copyright (2008) with permission from Elsevier.

injection cycles. Nevertheless, the property retention after three cycles under extreme recycling conditions (100% recycling with no added virgin polymer) indicated the promising recyclability of these materials.

24.6.3

Recycling of Polymer Nanocomposites

Significant improvements of various composite performances can be achieved using low or very low percentages of nanoparticles (1 to 10% wt, according to the types of nanoparticles), in comparison with usual micronic fillers [80, 81]. Lower composite densities can also be obtained, leading to energy savings. The progressive decrease of the production costs of nanoparticles has led to their accelerated diffusion and to a strong development of industrial formulations combining these emerging fillers with diverse additives and usual fillers or reinforcements. Nevertheless, toxicological and environmental concerns could put a break on the growth of nanoparticles. Potential risks linked to the production of nanoparticles, the transformation of nanocomposites, as well as their end of life, possibly accidental (case of fire) have been pointed out. This has recently led to address the end-of-life of this class of materials, regarding these potential risks, as well as the integration of these materials into life-cycle analysis and sustainable development approaches. In consequence, the possibility to recycle nanocomposites has to be considered globally, but also by focusing on the potential performances of the regenerated materials that can be issued from nanocomposites. The nanoparticles most frequently used in experimental works and industrial applications are lamellar silicates, able to exfoliate into polymers or to lead to the intercalation of these ones between mineral sheets after ion exchange with organophilic ions. Montmorillonite is the most cited lamellar silicate and the selection of the organophilic modifier allows its performances and the range of host polymers to be extended. In the case of organo-modified montmorillonite using alkyl ammonium ions, the influence of such modifiers on composite degradation during nanocomposites reprocessing has to be scrutinized in order to assess the recyclability of such emerging materials. Several processes could happen in function of the processing parameters and of the nature of the polymer. At first, a too high processing temperature can lead to a partial

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decomposition of the modifier by a Hoffman reaction, leading to a loss of the organophilic character of the layered silicate, which could entail a loss of the dispersion at the nanometric scale. Moreover, the degradation products of ammonium salts could also degrade the polymer during reprocessing if some chemical bonds of its structure are sensitive to ammonia. However, the research works dealing with the thermal degradation of nanocomposites are mainly devoted to the improvement of the thermal stability or fire reaction of virgin polymers in presence of nanoparticles. Few works concern the specific degradation effects caused by nanoparticles. The review made by Pandey et al. [82] has shown that organo-modified clays could contribute to degradation mechanisms by hydrolysis during reprocessing for polymers sensitive to this degradation mode (e.g. Polyamide 6). Xu et al. [83] have incorporated organo-modified montmorillonite using several modifiers in PET. These authors have shown that the extent of PET degradation, determined by molar mass measurements, depends on the surface chemistry of organo-modified clays. Nanocomposites based on clays presenting the highest rate in hydroxyl groups, exhibit the highest degradation rate. The percentage of alkylammonium salt is also presented as a critical parameter owing to catalytic degradation effects affecting PET and connected to the own degradation of organo-modifiers following the Hoffmann reaction. The possible decomposition of quaternary ammonium salts during the regeneration process of nanocomposites of engineering polymers such as polyesters, polyamides, and styrenic copolymers, could also limit the compatibility of clays with the host polymer for long residence times in internal mixers or extruders, particularly if successive reprocessing steps would be carried out. Nevertheless, it has been shown that pristine organo-modified clays could compatibilize polymer blends. This property could be advantageously applied to blend of reclaimed polymers. Sinha Ray et al. [84] have used organo-modified montmorillonite to compatibilize commercial polycarbonate/polymethyl methacrylate blends. Unfortunately, for HDPE/PA 6, the use of organo-modified clays could not lead to the expected mechanical properties. The inhibition of HDPE crystallization induced by the particles and the thermal degradation of modifiers were put forward by the authors to explain the unsatisfactory results. Nevertheless, a positive role of clay particles on blend microstructure highlighted the compatibilizing effect. The role of nanoparticle dispersion on PET degradation was also mentioned by Xu et al. [83]. A better dispersion increases the extent of interface between nanoparticles and polymer and, hence, the degradation processes involving the clay/PET interface. The use of chain extenders such as the pyromellitic dianhydride during the extrusion process has been proposed as solution. In addition, Swoboda et al. [54] have investigated the incorporation of submicronic kaolinite in recycled PET and PET/PC alloys. They have shown that the functionalization of mineral particles using a triphenylphosphite compound, acting as chain extender, could compensate for the adverse effect on thermal stability of the hydroxyl groups located on the kaolinite surface. In polyolefins, some authors [85] have also shown that the presence of organo-modified clays could decrease the temperature of elimination of side groups of ethylene vinyl acetate (EVA) in EVA/PE blends. Another critical process concerns the lower stability of nanocomposites regarding the photochemical aging in comparison with unfilled polymers. This phenomenon has been shown by Mailhot et al. [86], Morlat et al. [87] in PP and PE nanocomposites. It has been proved using infra red spectroscopy that the photochemical degradation is faster in presence of organo-modified clays. In addition, the induction period of the aging process is shortened. Among the authors who have really studied the behavior of nanocomposites after reprocessing, Thomson and Yeung [88] on the one hand, and Goitisolo et al. [89], on the other, have developed a standard methodology [50, 90], consisting in performing successive reprocessings of a composition of commercial or reclaimed polymer, blend or particulate filled composite in a polymer processing device (mono or twin screw extruder, injection molding) and to study the evolution of mechanical, thermomechanical properties and morphology of resulting materials.

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Figure 24.10 TEM micrographs at 60K magnification of the TPO nanocomposite for (a) Pass 1, (b) Pass 4, (c) Pass 7, and (d) Pass 10. Reprinted from [88]. Copyright (2006) with permission from Elsevier.

Thomson and Yeung [88] have studied successive re-extrusions using a twin screw extruder of a polyolefin elastomer (TPO) in presence of 9 %wt of PPgMA and 3 wt% of organo-modified clay, containing a quaternary ammonium salt. The compositions have been reprocessed up to 9 passes in the extruder. It has been shown that the increase of reprocessing steps entailed a better exfoliation of silicate particles. The authors ascribe the degradation of the polymer rather to the presence of the PPgMA compatibilizing agent than to the action of nanoparticles. Transmission Electron Microscope observations (Figure 24.10) revealed the presence of tactoids, suggesting the existence of a percolation threshold, able to determine the evolution of the rheological properties. This percolation effect would depend on the formation of carbonyl groups along the polymer chains and caused by their thermo-oxidation. Mechanical results highlighted a progressive decrease of properties with the number of reprocessing steps, except for the Yield stress, which remained stable. Thus, flexural and tensile modulus decreased for 10 passes, whereas the elongation at break fell from 15 to 9%. Except for elongation at break, all the properties have remained significantly higher than those of polymers without

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Figure 24.11 Elongation at break of nanocomposites in function of the number of passings, commercial nanocomposite (circles), made at laboratory (squares). Reprinted from [89]. Copyright (2008) with permission from Elsevier.

nanoparticles and compatibilizing agent. In order to limit the degradation processes, the authors suggested a combination of stabilizer and compatibilizer be added during the reprocessing steps. The study of the properties after successive retransformations by twin screw extruder of an organo-modified clay/PA 6 nanocomposite was investigated by Goitisolo et al. [89]. Five passages were made with two kinds of nanocomposites: a commercial one realized by polymerization in situ, the other made by mixing at molten state with organo-modified clay. The evolutions of tensile modulus and elongation at break were studied, as well as these of morphology and crystallinity. The degradation of the nanocomposite has been followed by the evolution of the color of material. The tensile modulus was not modified since the properties at low deformations are not generally damaged by moderate processes of thermochemical ageing. Nevertheless, the elongation at break strongly decreased after five passes, and more dramatically for the prepared nanocomposites, than for the commercial one (Figure 24.11). This seemed to prove the progressive degradation of the regenerated material. However, the degradation process of PA6 after successive re-extrusions was also noticed without nanoparticles. Consequently, it is not possible to conclude to a specific contribution of these ones, regarding thermo-oxidation or chain breakings. A complementary approach aiming to show the interest of nanoparticles in relation to maximize the performances of unfilled regenerated materials, has consisted in incorporating nanoparticles during reprocessing. Pegoretti et al. [91] have shown that the incorporation of 1 to 5 wt% of organo-modified or nonmodified lamellar silicate in a recycled PET, could entail an increase in tensile modulus up to 30%. This last value was obtained with the organo-modified one, which was only able to produce a nanometric microstructure. An improvement of the creep behavior of the filled regenerated materials was also noticed, the composition with the organo-modified clay at the highest rate led to the more advantageous values.

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24.7 Waste Management of Glass Fiber-reinforced Thermoset Plastics 24.7.1

Waste Management and International Context

Waste composite based on thermoset materials are originated from [92]: 1. The pre-consumption materials or production waste of rejected pieces (estimated at 3% of 112.000 t in a year in Europe). 2. The post-consumption materials or end-of-life waste such as the End-of-Life Vehicle (ELV), Waste Electric Electronic Equipment (WEEE), boats. . . Thermoset polymer composites are generally perceived as difficult or impossible to be recycled, conversely to thermoplastics, which can be remelted. Nevertheless, some technologies were developed to reprocess composite waste and to avoid landfilling. Methods to treat composite waste must be able to adapt to the characteristics of these multicomponent materials, with a minor fraction of polymer matrix (often lesser than 30% in weight in the case of SMC), and often contaminated in the case of end-life waste. Regeneration processes have to take into this difficulty [13]. Figure 24.12 presents a synthesis of the various routes of composites waste management developed or explored. The main industrial producers of composite materials based on thermoset polymers have tried to solve the problem end-of-life materials. Because of own regulations and different cultures, the proposed solutions differ, but the problems remain the same. Within this framework, many programs and studies were undertaken since the beginning of the nineties in order to find ways of regeneration or waste management by mechanical recycling, chemical degradation, pyrolysis, combustion, etc. [92]. In Europe, waste management of end-of-life composites is a main concern of the European Group of Composite Producers (GPRMC), which became EuCIA in June 2004 (European Composites Industries Association). EuCIA considers that the recyclability is essential for the future development of the composites, at the same level as health and safety. New prohibitions are considered as regards the landfilling of end-to life products and even the incineration although it can have the merit to recover energy [92]. More and more of the final users facing to find waste elimination solutions have decided to integrate recyclability in new production processes (93). Another objective of EuCIA is to diffuse knowledge on the recyclability of materials, in particular while referring to the activities of recycling installation by companies such as ERCOM Composites Recycling in Germany, Mixte Composite Recycling in France, Lonza in Italy and Miljotek in Norway. These installations are mainly dedicated to the treatment of the Sheet Molding Compounds (SMC – unsaturated polyester reinforced by glass fibers and loaded by calcium carbonate). In addition, a project aiming to promote new solutions of end-of-life for the recycling and the reuse of the composites was launched in August 2003 by the main European suppliers of composites. The participants could affix on their products the Green FRP recycling label. In the 1990s, in North America, the recycling of thermosets and thermoset composites problem was the subject of many studies, in particular about the treatment of composite waste SMC by crushing, carried out by the Phœnix Fiberglass Inc. company (Canada), which filed for bankruptcy in 2007, since recycled goods were produced at too high a price. The Seawolf Industries Inc. company developed a process of crushing making it possible to guarantee the integrity of glass fiber. Taking into account the capital cost of expensive recycling, tests of reincorporation of SMC/BMC were studied by various companies (SMC Automotive Alliance (SMCAA) or Composite Corp Technologies. (CTC)), regrouping within corporation of companies. The latter developed snowboards and basket panels starting from industrial offcuts by extrusion/compression with low pressure of recycled SMC materials mixed with glass fibers and incorporated in thermoplastics.

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Matter valorization Pulverization

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Figure 24.12 Various techniques of treatment and valorization of composite materials-based polyester thermosets. Reprinted from [124]. Copyright (2002) Savoie Technolac.

In Japan, the question of composites waste processing is relatively long-standing. In 1974, the JPRS (Japan Reinforced Society Plastics) created a committee of research about waste processing, which was primarily interested in composite waste processing. In 1991, a regulation about recycling came into effect. Finally, in 1998, the JRPS created the centre of promotion for the recycling of composites [94].

24.7.2

Feedstock Recycling by Pyrolysis

Feedstock recycling techniques are based on a non-oxidative heating leading to the scission of a polymer in small molecules. During the decomposition of thermoset composites based on polyester and epoxy resins, three main products of pyrolysis are formed: pyrolysis, gases, oils and solid residues [95, 96]. Pyrolysis leads mainly to olefins and other hydrocarbons which can be used as fuels or raw materials in petrochemistry. In 1990, pyrolyzed composite wastes based on thermoset resins were successfully incorporated in tires [95]. Moreover, the gases and oils recovered could be used for their energy values, of the same as for natural gas [97], while the residual solid products could also be incorporated into concrete [98]. The SMC Automotive Alliance has also developed a pyrolysis process of automobile waste leading to 70% solid residues and oil and 30% gas. The solid phase was crushed and milled to serve as fillers in thermoplastics [99, 100]. Thus, as

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has been mentioned above, the combined use of pyrolysis and crushing has been proved to be an interesting recycling global process, from the technical and economical point of view [101]. 24.7.3

Solvolysis or Chemical Recycling

The chimiolysis of thermoset resins corresponds to the use of reactive solvents causing the breaking of the ester linkages and the destruction of the three-dimensional network. It is then possible to separate the polymeric matrix from fibers and the fillers. The chimiolysis generally used for composites is hydrolysis or alcoolysis in neutral or basic environment. Other methods, in particular aminolysis, are investigated for SMC. Hydrolysis: Winter [102] has described the degradation of crushed SMC or BMC (Bulk Molding Compound) in the presence of ethanol/water (2:1) with 85◦ C, catalyzed by potash. After 48 h, the conversion rate was 86%. However, this process presents several disadvantages, since it cannot regenerate either glass fibers or the fillers:

r r r

glass fibers were degraded by potash a stage of neutralization to the acid was necessary, leading to the formation of salts which required costs of important treatments calcium carbonate was destroyed owing to neutralization.

Glycolysis: According to Yoon et al. [103], it has been possible to make crushed polyester resins (Ø < 500 μm) by reaction, initially, with the propylene glycol in excess without catalysts at 170◦ C during 3 h and then 4 h at 225◦ C. Then, in a second process, a reaction of polycondensation with maleic anhydride (24 h at 150◦ C) was carried out and finally, after a distillation of styrene, this led to a new unsaturated polyester resin with properties considered to be acceptable. A patent of Ashland Oil [104] has described a method of glycolysis of crushed SMC (Ø < 1000 μm) using diethylene glycol or dipropylene glycol, allowing filler and glass fibers to be separated. The fraction rich in matrix underwent a glycolysis at a temperature ranging between 220 and 270◦ C with or without catalyst of transesterification (sodium methoxyde or potash). The reaction was complete after 30 h. The resulting polyols could react with polyisocyanates to make polyurethanes or with unsaturated carboxylic acids to make unsaturated polyesters. The Miyaso Chemical Company, through two patents of glycolysis [105, 106], described a process of recycling of micronized thermosets using ethylene or di- or tri-ethylene glycol with a conversion rate of respectively 62% for 3 h at 270◦ C for the first one and 80% in 5 h with 250◦ C for the second. The glycols obtained were also reused to make unsaturated polyesters or polyurethanes. Aminolysis: Funke et al. [107] have described aminolysis of the thermoset polyester resins using benzylamine catalyzed by ammonium chloride. The authors noticed that. after 3 h, all ester bonds were broken. More recently the work of Tai et al. [108] described the decomposition of end-of-life polyesters resins micronized using diaminated compounds as xylylene diamine. This solvolysis led to a total digestion of the resin after 2 h and at only 170◦ C. Moreover, liquid products were obtained. Winter et al. [102] proposed the use of reagents containing aminoalcool such as monoethanol amine (MEA). A separation of the three principal fractions of the SMC (matrix (soluble in methanol)/filler/fibers) could be obtained after 48 h. MEA reacted with the amine function to give imides and to the rupture of the ester bonds. The fibers could be reused up to a loading of 50% in new composites like BMC. In the same way, work carried out by Vall´ee et al. [109] and Perrin et al. [110] proposed two modes of aminolysis of the SMC crushed beforehand (Ø < 1000 μm), by attack of a triamine (diethylenetriamine

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Figure 24.13 Mechanism of aminolysis derived from the reaction between DETA and UP (Et: CH2-CH2). Reprinted from [109]. Copyright (2004) with permission from American Chemical Society.

DETA) or a diamine (xylylene diamine XDA), on the esters bonds of the thermoset resins. A complete digestion of the resin occurs after 10 h at temperatures from 200◦ C to 250◦ C. The mechanisms of the reactions of solvolysis are indicated Figure 24.13 and led to the formation of imides. Other routes of solvolysis: Other processes can be mentioned, such as the ozonolysis under pressure which allows the polymer network to be partially destroyed and the fibers to be partially separated [111]. Another process is based on the catalysis by lipases and a polar solvent, this last causes swelling, and can solubilize the degradation products [112].

24.7.4

Mechanical Recycling of Glass-reinforced Thermoset Composites

Another possible route of regeneration is the mechanical recycling of materials after size reduction. The composites are currently reduced in two stages: shredding (of pieces of a few square centimeters) and crushing. Primary grinding (or shredding) makes it possible to obtain parts adapted to a secondary treatment (e.g. addition of interfacial agents). Crushing allows micronized powder or fibers to be obtained, which are then separated and collected according to their particle size distribution. A separation of the metal inserts is

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generally required [92]. The resulting micronized fraction can replace usual mineral fillers in polymers. The fibers are separated beforehand for reuse as reinforcement as short fibers. In a patent from the Phoenix Fiberglass company, fibers were recovered while remaining waste was micronized as powders [92]. Seawolf Industries Inc. also developed a crushing process, making it possible to recover fibers without destroying them, the recycled material being usable as reinforcement of composites [92]. The final product was mixed with a synthetic foam containing polyester which can be pulverized, named Scrap Recycling System, and intended for the production of articles such as bathtubs or boat decks. Various applications have been developed for micronized powder and fibers. The use of formulations containing powders used as fillers requires a thorough control of viscosity and composition. Moreover, the mechanical characteristics of blends containing powder and recycled glass fibers can be lowered in comparison with the corresponding usual commercial compositions. This often leads to open loop applications for which mechanical characteristics are not dominant issues [92]. Regenerated glass fibers or fibrous fractions issued from mechanical recycling of thermoset composites have been used as reinforcement agents in asphalt or concrete as well as in thermoplastic matrices. Lastly, Plastic Omnium Company (113), in partnership with MCR/Inoplast, has incorporated a mixture of regenerated glass fibers resulting from SMC and new glass fibers in a polypropylene matrix. These compositions were developed for under-the-hood applications in Renault vehicles. In Europe, since the 1990s two important facilities have been dedicated to the mechanical recycling of end-of-life composite materials such as SMC/BMC and epoxy composites: ERCOM in Germany and MCR in France. MCR produces three categories of products [113]:

r r r

fine powder (lower than 300 μm) short fibers (around 1 mm) long fibers (around 10 mm).

The powder is recycled as fillers in thermoset composites compositions but its price is expensive compared to calcium carbonate. It is also usable in concretes. Short fibers are partially recycled in thermoplastics, whereas long fibers are incorporated in asphalt, concretes or SMC/BMC composites [114, 115]. From an economic point of view, the selective reuse of the fibrous fraction seems much more promising than the recycling of the powder since the value of reinforcements is higher. The fraction of recovered fibers can be reused successfully in thermoset composites, in combination with virgin fibers, without affecting seriously the properties of the material, in a proportion up to 50 wt% of the global reinforcement, and with a cost saving which can exceed 30%. Nevertheless, a reuse at 100% of the regenerated fibrous fraction incorporated into new thermoset composites involves a degradation of the mechanical properties. A very promising alternative solution is the reuse of the fraction of recovered fibers as reinforcement of thermoplastic matrices, such as polypropylene and polyethylene [13]. 24.7.4.1

Reuse in Thermoset Polymers

Epoxy Resins Investigations carried out at the Clean Washington Center (USA) indicate that there exists a ‘profitable’ way of treatment as a reinforcement of the ground offcuts in SMC [116]. Shredding and micronization by hammer mill have been made in order to obtain average lengths of fibrous particles of 2 mm. The incorporation of this reinforcing fraction showed that only 1 %wt of the recycled fibers could be added to the epoxy resin until viscosity reaches unacceptable levels. Fiber length affects viscosity, and particularly the longest fibers. The results also showed that the mechanical strength of the epoxy resin increased by 16% with the addition of 1% fiber having a length lower than 0.5 mm. These ones entailed not

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only a higher resistance compared with fibers of 0.5 mm, but also reduced the viscosity, at constant fiber loading. Lastly, the University of Michigan developed the use of fillers from crushed and micronized polyester composites based on cutoffs to make low cost prototype epoxy molds, in order to replace thermoforming virgin resin up to 70 %wt [104]. However, some complementary studies are necessary to determine the commercial feasibility of this technique on a larger scale. Polyester Resin A study performed at CEER/University Alfred (USA) [117], tried to investigate the structure–properties relationship of a composite made of a ground SMC coming from the MCR company, incorporated in a molded polyester resin. It was reported that the reduction in mechanical strength of the resulting composites in comparison with a commercial one at the same loading, was due to a lesser amount of fibers. As a matter of fact, it regenerated fibers are surrounded by a thick coating of residual polyester. Moreover, this reduction of mechanical strength was also attributed to the weak adhesion between the residual resin remaining at the fiber surface and the virgin resin. Surface treatments of regenerated fibers could be proposed to improve adhesion. Pedroso et al. [118] reintroduced offcuts of composites-based polyester matrix (shredded, crushed and sieved) coming from public phone boxes in a virgin polyester resin. Mechanical properties as well as the surface aspect of these recycled composites were considered. This work showed that the influence of a slightly viscous resin could maximize the mechanical properties. A research group from the Technical University of Athens [13] has studied the size reduction of thermoset materials into fine powders. Reinforced glass cutoffs of BMC were pulverized and the powder obtained (sieved at 74 μm) was used as filler in SMC at various percentages, from 0 to 88 wt%. A microscopic examination of the fillers indicated that they were made up of glass fibers as well as rough spherical particles. A mechanical test indicated that the inclusion of crushed BMC in SMC led to a reduction in the mechanical properties, particularly for the highest filler percentages.

24.7.4.2

Reuse in Thermoplastic Polymers

Investigations have been carried out at Mines Al`es (France) by Reygrobellet et al. [119] and more recently by Perrin et al. [110, 120, 121] on treatments of fibrous fractions of ground SMC, in order to reincorpore it in thermoplastic polymers. An original process of crushing/sieving and surface modification, called Syltec, was developed to recover and to reuse short fibers of thermoset composites waste as reinforcement. The maximization of the mechanical properties of thermoplastic composites reinforced with these fractions containing calcium carbonate, E-glass fibers and polyester matrix, was aimed at, in order to compete with thermoplastic composites reinforced by virgin glass fibers. Syltec is composed of four stages. The first stage, phase of conditioning, intends to separate glass fibers from the polyester resin and to calibrate the size of this fraction. The use of a shredder allows the length of fibers to be maintained and the global composition of the reinforcement to be kept constant, particularly the fibrous content. The second stage consists in a controlled dissolution of calcium carbonate present in the reinforcement by an acid-base reaction and leading to a reinforcing fraction of 40–45% in mass. In the third stage, ionic interactions are created: on the one hand between calcium carbonate and the anhydride of the thermally activated polyester, by functionalization of polyester; on the other hand, between calcium carbonate and the anhydride function of PPgma , acting as interfacial agent. Hence, a reinforcement of composite SMC/PP matrix is achieved, allowing the mechanical performance to be increased. The last stage, only carried out in the case of a nonpolar host polymer, corresponds to a chemical functionalization between the modified regenerated glass fibers and the polymer matrix (PP/PPgma ), using a

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6000 5000

PP matrix with SMC composite filler (First stage of grinding)

4000 3000

PP matrix with SMC composite filler (Second stage of partial dissolution)

2000

PP matrix with SMC composite filler (Third stage of thermal treatment)

1000 4

0

8 12

10

16

20

20

30

24

40 50 60

Rupture energy (kJ/m2)

Yield stress (MPa)

Young’s modulus (MPa)

6000 5000 Reference (raw PP reinforced by coated glass fibers

4000 3000

PP matrix with SMC composite filler (Third stage of thermal treatment)

2000

PP matrix with SMC composite filler (Fourth stage of chemical treatment)

1000 4 8 12 16 20 24

0 10 20 30 40 50

Rupture energy (kJ/m2)

60

Yield stress (MPa)

Figure 24.14 Evolution of mechanical properties of ground SMC into PP matrix regarding the stage of Syltec process. Reprinted from [119].

polyfunctional agent (e.g. an amine). Condensation in imide function thus allowed the creation of a chemical bonded interface whereas it was only ionic after the third stage. Comparable mechanical properties with those of commercial sized glass fibers incorporated in the same PP matrix have been achieved using this regenerated reinforcement (Figure 24.14) [122]. Moreover, this process was adopted successfully on polar matrices (PA-6) and can be transposed to other polyester resins. The Syltec process has been scaled up and obtained the Prize of the Innovative Techniques in 2006, awarded by the French Agency of Environment (ADEME).

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Similar Process with Carbon Fiber-reinforced Epoxy Resin To end this chapter with a correlation between a recycling process from glass fibers and other kind of fiber-reinforcing thermosets by incorporation in thermoplastics, McNally et al. [123] have used regenerated carbon fibers in high density polyethylene from an epoxy/CF composite used in Formula 1 motor racing cars. Fibers were reclaimed using the CARBONCLEAN technology, a thermochemical process that recovers 95% of the CF. The unsized CF obtained were subsequently milled to produce fibers with a density of 1.8 g.cm−3 , a nominal length of 250 μm and a diameter of 7 μm. Both the Young’s modulus and ultimate tensile stress increased with increasing CF loading, but the percentage stress at break was unchanged up to 5 wt% loading, then decreased with further successive addition of CF. The electrical conductivity of the PE matrix was enhanced by about 11 orders of magnitude on addition of recycled CF with a percolation threshold of 7 and 15 wt%.

24.8 Conclusion Recycling of filled and fiber-reinforced polymer is one of the main challenges for the plastic composite industry at present time. Significant advances have been made in the field of sorting and identification of blended reclaimed plastics. Nevertheless, the increase in purity of regenerated materials through recycling processes requires the development of combined techniques, able to detect the presence of fillers, reinforcements and additives. Many techniques are available to separate and regenerate the components of multiphase polymers systems containing fillers or fibers, such as resin dissolution, mechanical, chemical and thermal processes. More recently, according to the requirements of new regulations, new methods have been developed to extract pigments and additives, particularly halogenated flame retardants, progressively phased out in electric and electronic equipments in Europe and other developed countries. Mechanical recycling of thermoplastic composites performed on pure reclaimed materials can lead to high level of mechanical performance for the recycled polymers, provided that no significant degradation such as hydrolysis occurs during reprocessing. It has been shown that the presence of fillers in the waste polymers could act diversely on the stabilization during and after reprocessing. Reactive surface treatments performed on the functional groups of the filler surface could prevent the negative effect of their surface chemistry on the polymer stability during reprocessing. Hence, grafted polymers acting as compatibilizers can strongly improve the performances of recycled composites. They can notably limit the degradation of aspect ratio of fibers which strongly improves the mechanical properties of reinforced thermoplastics. As a matter of fact, fillers having a high aspect ratio allow the tensile mechanical properties of recycled polymers to be upgraded. This effect is particularly noticeable for nanoparticles such as organo-modified clays. On the whole, the intrinsic effect of nanoparticles on polymer degradation during reprocessing is not significant. Nevertheless, particular attention has to be focused on the presence of impurities and on the possible decomposition of organo-modifier, which can lead to the degradation of the host polymer. Compared to thermoplastic waste, thermoset composites materials have been considered difficult to recycle up to now. The industry is always facing an environmental and economic challenge, even if various projects and experiences in pilot and medium-scale plants have demonstrated that mechanical recycling is technically possible. Research works carried out on the reincorporation of selected and modified fractions of ground SMC composites in thermoplastics have shown that high levels of performances can be achieved in comparison with the same polymers reinforced with pristine glass fibers. Polymer materials are examples of a fabricated construction that does not have a cycle. These items do not deteriorate and they merely fill up landfills. Plastic needs help to be recycled. Nevertheless, one has to recognize that some filled plastics are very difficult to recycle. Many complications are due to the presence of specific additives, such as different flame retardant or nanoparticles in new

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generations of composites, which have to be well-identified and well-sorted before being reused. Some new techniques are being developed such as the on-line (1 t/h rate) Near InfraRed (NIR) technique for the identification of thermoplastics – mainly Waste of Electric and Electronic Equipments – or X-fluorescence for the sorting of fillers. The improvement in these techniques will be a main concern for the management of plastic wastes in the near future.

References 1. Composite recycling in the construction industry, A. Conroy, S. Halliwell, T. Reynolds, Composites: Part A, 37, pp 1216–1222 (2006). 2. Environmental impacts by disposal of plastic from municipal solid waste, C. Molgaard, Resources, Conservation and Recycling, vol. 15, pp 51–63 (1995). 3. ECRC heads search for composites recycling solutions, D. McKechnie, T. Wegman, Reinforced plastics (July 2007). 4. Recycling collaborative combats legislation threat, G. Marsh, Reinforced plastics (September 2005). 5. Recycling threat to Europe’s composites industry, Reinforced plastics (September 2006). 6. Recycling of plastics: Automatic identification of polymers by spectroscopic methods, J. Florestan, A. Lachambre, N. Mermilliod, J.C. Boulou, C. Marsifi, Resources, Conservation and Recycling, vol. 10, pp 67–74 (1994). 7. Plastic separation planning for end-of-life electronics, J. Williams, E. Grant, P. Rios, L. Blyler, L. Tieman, L. Twining, W. Bonawi-tan, M. Madden, N. Guthrie, IEEE Transactions on Electronics Packaging Manufacturing, vol. 29(n◦ 2), pp 110–118 (2006). 8. The use of electrostatic techniques for the identification and sorting of waste packaging materials, G. Hearn, J. Ballard, Resources, Conservation and Recycling, vol. 44, pp 91–98, (2005). 9. Electrostatic separation of plastics by friction mixer with rotary blades, Y. Matsushita, N. Mori, T. Sometani, Electrical Engineering in Japan, vol. 127(n◦ 3), pp 33–40 (1999). 10. The triboelectric pen: an electrostatic method for the identification of plastics in recycling, G. Hearn, P. Mucci, A. Eyres, IAS’96: IEEE Industry Applications Conference (October 1996). 11. Chemical recycling of carbon fiber composites using alcohols under subcritical and supercritical conditions, R. Pinero-Hernanz, J. Garcia-Serna, C. Dodds, J. Hyde, M. Poliakoff, M. Cocero, S. kingman, S. Pickering, E. Lester, The Journal of Supercritical Fluids, vol. 46, pp 83–92 (2008). 12. Recycling of the products obtained in the pyrolysis of fiber-glass polyester SMC, I. de Marco, J. Legarreta, M. Laresgoiti, A. Torres, J. Cambra, M. Chomon, B. Caballero, Journal of Chemical Technology and Biotechnology, vol 69, pp 187–192 (1997). 13. Recycling of the fibrous fraction of reinforced thermoset composites, C. Kouparitsas, C. Kartalis, P. Varelidis, C. Tsenoglou, C. Papaspyrides, Polymer Composites, vol. 23(n◦ 4), pp 682–689 (2002). 14. Recycling technologies for thermoset composite materials – current status, S. Pickering, Composites: Part A, vol. 37, pp 1206–1215 (2006). 15. Recycling of composite materials, M. Buggy, L. Farragher, W. Madden, Journal of Materials Processing Technology, vol. 55, pp 448–456 (1995). 16. Recycling of glass fiber reinforced thermo-plastic composites. Part I. Ionomer and low density polyethylene based composites, C. Papaspyrides, J. Poulakis, C. Arvanitopoulos, Resources, Conservation and Recycling, vol. 14, pp 91–101 (1995). 17. Recycling of polypropylene-based composites, J. Poulakis, P. Varelidis, C. Papaspyrides, Advances in Polymer Technology, vol. 16(n◦ 4), pp 313–322 (1997). 18. Characterization of LDPE residual matrix deposited on glass fibers by a dissolution/reprecipitation recycling process, N. Zafeiropoulos, P. Varelidis, C. Papaspyrides, T. Stern, G. Marom, Composites: Part A, vol. 30, pp 831–838 (1999). 19. Separation of a thermotropic liquid crystalline polymer from polypropylene composites, M. Collier, D. Baird, Polymer Composites, vol. 20(n◦ 3), pp 423–435 (1999).

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20. Extraction of components of composite materials: ionic liquids in the extraction of flame retardants from plastics, H. Lateef, S. Grimes, R. Morton, L. Mehta, Journal of Chemical Technology and Biotechnology, vol. 83, pp 541–545 (2008). 21. Extraction of brominated flame retardants from polymeric waste material using different solvents and supercritical carbon dioxide, A. Altwaiq, M. Wolf, R. van Eldik, Analytica Chimica Acta, vol. 491, pp 111–123 (2003). 22. Thermally crosslinkable-decrosslinkable system using diepoxy crosslinkers containing sulfonate ester moiety, Y-D. Shin, A. Kawaue, H. Okamura, M. Shirai, Reactive & Functional Polymers, vol. 61, pp 293–302 (2004). 23. Characterization of thermally reworkable thermosets: materials for environmentally friendly processing and reuse, J-S. Chen, C. Ober, M. Poliks, Polymer, vol. 43, pp 131–139 (2002). 24. A new concept for recycling of thermosetting resins I: The case of crosslinkable copolyesters, J. Economy, A. Andreopoulos, Polymers for Advanced Technologies, vol. 7, pp 561–570 (1996). 25. Chemical recycling of poly(ethylene terephtalate), G. Karayannidis, D. Achilias, Macromolecular Materials and Engineering, vol. 292, pp 128–146 (2007). 26. Reactive & functional polymers, M. Modesti, F. Simioni, R. Munari, N. Baldoin, vol. 26, pp 157–165 (1995). 27. (Supercritical) ammonia for recycling of thermoset polymers, W. Mormann, P. Frank, Macromolecular Symposia, vol. 242, pp 165–173 (2006). 28. Hydrogenolytic degradation of thermosets, D. Braun, W. von Gentzkow, A. Rudolf, Polymer Degradation and Stability, vol. 74, pp 25–32 (2001). 29. Mechanical characteristics of modified unsaturated polyester resins derived from poly(ethylene terephtalate) waste, M. Farahat, Polymer International, vol. 51, pp 183–189 (2002). 30. Desaminated glycolysis of water-blown rigid polyurethane foams, J. Lee, D. Kim, Journal of Applied Polymer Science, vol. 77, pp 2646–2656 (2000). 31. Supercritical secondary alcohols as useful media to convert polyamide into monomeric lactams, A. Kamimura, Y. Oishi, K. Kaiso, T. Sugimoto, K. Kashiwagi, ChemSusChem, vol. 1, pp 82–84 (2008). 32. Chemical recycling of carbon fiber reinforced composites in nearcritical and supercritical water, R. Pinero-Hernanz, C. Dodds, J. Hyde, J. Garcia-Serna, M. Poliakoff, E. Lester, M. Cocero, S. Kingman, S. Pickering, K. Wong, Composites: Part A, vol. 39, pp 454–461 (2008). 33. Poly(ethylene terephtalate) glycolysis under microwave irradiation, A. Krzan, Polymers for Advanced Technologies, vol. 10, pp 603–606 (1999). 34. Microwave-assisted depolymerization of poly(ethyelen terephtalate) [PET] at atmospheric pressure, M. Alavi Nikje, F. Nazari, Advances in Polymer Technology, vol. 25(n◦ 4), pp 242–246 (2006). 35. The characterization and reuse of glass fibers recycled from scrap composites by the action of a fluidized bed process, J. Kennerley, R. Kelly, N. Fenwick, S. Pickering, C. Rudd, Composites Part A, vol. 29A, pp 839–845 (1998). 36. A fluidized-bed process for the recovery of glass fibers from scrap thermoset composites, S. Pickering, R. Kelly, J. Kennerley, C. Rudd, N. Fenwick, Composites Science and Technology, vol. 60, pp 509–523 (2000). 37. The properties of glass fibers recycled from the thermal processing of scrap thermoset composites, J. Kennerley, N. Fenwick, S. Pickering, C. Rudd, Journal of Vinyl & Additive Technology, vol. 3(n◦ 1), pp 58–64 (1997). 38. Tertiary recycling of automotive plastics and composites, R. Allred, L. Busselle, Journal of Thermoplastic Composite Materials, vol. 13, pp 92–101 (2000). 39. Recycling of fiber-reinforced polymeric waste by pyrolysis: thermo-gravimetric and bench-scale investigations, A. Cunliffe, N. Jones, P. Williams, Journal of Analytical and Applied Pyrolysis, vol. 70, pp 315–338 (2003). 40. Characterisation of products from the recycling of glass fiber reinforced polyester waste by pyrolysis, A. Cunliffe, P. Williams, Fuel, vol. 82, pp 2223–2230 (2003). 41. Recycling of pyrolysis of thermoset composites: characteristics of the liquid and gaseous fuels obtained, A. Torres, I. de Marco, B. Caballero, M. Laresgoiti, J. Legarreta, M. Cabrero, A. Gonzalez, M. Chomon, K. Gondra, Fuel, vol. 79, pp 897–902 (2000). 42. GC-MS analysis of the liquid products obtained in the pyrolysis of fiber-glass polyester sheet moulding compound, A. Torres, I. de Marco, B. Caballero, M. Laresgoiti, M. Cabrero, M. Chomon, Journal of Analytical and Applied Pyrolysis, vol. 58–59, pp 189–203 (2000).

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43. Monomer recovery from aluminium hydroxide high filled poly(methyl methacrylate) in a fluidized bed reactor, G. Grause, M. Predel, W. Kaminsky, Journal of Analytical and Applied Pyrolysis, vol. 75, pp 236–239 (2006). 44. Feedstock recycling of synthetic and natural rubber by pyrolysis in a fluidized bed, W. Kaminsky, C. Mennerich, Z. Zhang, Journal of Analytical and Applied Pyrolysis, vol. 85, pp 334–337 (2009) 45. Investigations on the recycling of hemp and sisal fiber reinforced polypropylene composites, A. Bourmaud, C. Baley, Polymer Degradation and Stability vol. 92, pp 1034–1045 (2007). 46. Polymer degradation and stabilization, N. Grassie, G. Scott, Cambridge University Press (1988). 47. Kinetics of hydrolysis and thermal degradation of polyester melts, K.S. Ceo, J.D. Cloyd, J. of Applied Polymer and Science, vol. 42, pp 845 (1991). 48. Chap. 4: recycling of PET, J. Scheirs, Polymer Recycling, John Wiley and Sons (1998). 49. Chain extension of recycled poly(ethylene terephtalate) with 2,2’-bis-(2-oxazoline), N. Cardi, R. Po, G. Gianotta, E. Occhiello, F. Garbassi, G. Messina, J. of Applied Polymer and Science, 50, pp 1501 (1993). 50. Recycling of high density polyethylene containers, M. Kostadinova Loultcheva, M. Proietto, N. Jilov, F. P. La Mantia, Polym. Degrad. Stab., vol. 57, pp 77–81 (1997). 51. Influence of the wood fiber filler on the internal recycling of poly(vinyl chloride)-based composites, L. Augier, G. Sperone, C. Vaca-Garcia, M.-E. Borredon, Polymer Degradation and Stability 921169–1176 (2007). 52. Ageing and stabilisation of filled polymers: an overview, N. S. Allen, O. M. Edge, T. Corrales, A. Childs, C. M. Liauw, F. Catalina, C. Peinado, A. Minihand D. Aldcroft, Polymer Degradation and Stability 61, 183–199 (1998). 53. Photochemical stabilization of linear low density polyethylene/clay nanocomposites : towards durable nanocomposites, S. Morlat-Therias, E. Fanton, J.L. Gardette, N. Tzankova Dincheva, F.P. La Mantia, V. Malatesta, Polymer Degradation and Stability, vol. 93, pp 1776–1780 (2008). 54. Organomodified ultrafine kaolin for mechanical reinforcement and improved flame retardancy of recycled polymer terephtalate, in Fire Retardancy of Polymers, B. Swoboda, E. Leroy, J.M. Lopez Cuesta, C. Artigo, C. Petter, C.H. Sampaio, Ed T.R. Hull, B. Kandola, RSC Publishing, Cambridge (2009). 55. Product design with glass fiber reinforced polymer blends with potential applications in recycling, M. Xanthos, Polymer Composites, vol. 19, pp 768–780 (1998). 56. Effect of recycling on the mechanical properties of wood fiber-polystyrene composites. Part I: Chemithermomechanical pulp as a reinforcing filler, D. Maldas, B. V. Kokta, Polymer Composites, vol. 11, pp 77–83 (1990). 57. Effect of ground rubber powder on properties of natural rubber, S. Li, J. Lamminm¨aki, K. Hanhi, Macromolecular Symposia, vol. 216, pp 209–216 (2004). 58. Reclamation and recycling of waste rubber, B. Adhikari, D. De, S. Maiti, Progress in Polymer Science, vol. 25, pp 909–948 (2000). 59. Recycled plastic compounds containing ground rubber tires as filler, P. K. Pramanik, B. Dickson, Antec’95, pp 3738–3742 (1995). 60. Surface activation of crosslinked rubbers, and especially ground tyre rubber, via monomer grafting, I. Fuhrmann, J. Karger-Kocsis, SPE-ANTEC, San-Francisco (2002). 61. Compatibilisation of polyethylene/ground tyre rubber blends by γ irradiation, R. Sonnier, E. Leroy, L. Clerc, A. Bergeret, J.M. Lopez-Cuesta, Polymer Degradation and Stability, vol. 91, pp 2375–2379 (2006). 62. Compatibilizing thermoplastic/ground tyre rubber powder blends: Efficiency and limits, R. Sonnier, E. Leroy, L. Clerc, A. Bergeret, J.-M. Lopez-Cuesta, A.-S. Bretelle, P. Ienny, Polymer Testing, Vol. 27, Issue 7, pp 901–907 (2008). 63. Methods for polyurethane and polyurethane composites, recycling and recovery: A review, K. M. Zia, H. N. Bhatti, I. A. Bhatti, Reactive and Functional Polymers, vol. 67, pp 675–692 (2007). 64. Chap. 12: environmental aspects of fillers, G. Wypich, in Handbook of Fillers, Chem Tec Publishing, New York (1999). 65. Further studies on the use of pulverised flexible polyurethane foam particles as fillers in natural rubber compounds, N. Sombatsompop, G.L.A Sims, Cellular Polymers, vol. 15, pp 317–334 (1996). 66. Stability parameters for mineral suspensions: Improving the dispersion of fillers in thermoplastics, A. Lafaurie, N. Azema, L. Ferry, J.-M. Lopez-Cuesta, Powder Technology, article online Sciencedirect (2009).

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67. Use of natural fibres and environmental aspects, H. Burger, A. Koine, R. Maron, K.P. Mieck, International Polymer Science Technology, vol. 22, pp 25–34 (1995). 68. The effect of recycling on the time-dependent behavior of polycarbonate reinforced with short glass fibers, A.D. Drozdov, A. Al-Mulla, R.K. Gupta, Composites Science and Technology 64, pp. 129–144 (2004). 69. Improvement of the mechanical properties of an HDPE/PS blend by compatibilization and incorporation of CaCO3, F. Sahnoune, J. M. Lopez Cuesta, A. Crespy, Polymer Engineering and Science, vol. 43, pp 647–660 (2004). 70. lnfluence of impurities on mechanical properties of recycled glass fiber reinforced polyamide 66, P.-A. Eriksson, A.-C. Albertsson, P. Boydell, J.A.E. Manson, Polymer Engineering and Science, pp. 749–756 (1998). 71. Prediction of mechanical properties of recycled fiberglass reinforced polyamide 66, P.-A. Eriksson, A.-C. Albertsson, P. Boydell, G. Prautzch, J.A.E. Manson, Polymer Composites, vol. 17, pp 830–839 (1996). 72. Analysis of the effect of mechanical recycling upon tensile strength of a short glass fiber reinforced polyamide 6,6, A. Bernasconi, D. Rossin, C. Armanni, Engineering Fracture Mechanics, vol. 74, pp 627–641 (2007). 73. Characterization of mixed fiber nylon composites incorporating composite scrap, A. Licea-Claverie, F. J. U. Carillo, A. Alvarez-Castillo, V.M. Castagno, Polymer Composites, vol. 20, pp 314–320 (1998). 74. Impact strength of recycled thermoplastic composites subjected to corrosive environment, G.C. Papanicolaou, D. Karagiannis, D.A. Bofilios, J.H. Van Lochem, C. Henriksen, H.H. Lund, Polymer Composites, vol. 29, pp 1026–1035 (2008). 75. Recyclability of a continuous E-glass fiber reinforced polycarbonate composite, J. Chu, J. L. Sullivan, Polymer Composites, vol. 17, pp 556–567 (1996). 76. Influence of the wood fiber filler on the internal recycling of poly(vinyl chloride)-based composites, L. Augier, G. Sperone, C. Vaca-Garcia, M.-E. Borredon, Polymer Degradation and Stability, vol. 92, pp 1169–1176 (2007). 77. Mechanical properties of short flax fiber bundle/polypropylene composites: Influence of matrix/fiber modification, fiber content, water uptake and recycling, A. Arbelaiz, B. Fernandez, J.A. Ramos, A. Retegi, R. Llano-Ponte, I. Mondragon, Composites Science and Technology, vol. 65, pp 1582–1592 (2005). 78. Processing of natural-fiber reinforced polymers and the resulting dynamic–mechanical properties, B. Wielage, Th. Lampke, H. Utschick, F. Soergel, Journal of Materials Processing Technology, vol. 139, pp 140–146 (2003). 79. Effect of recycling on mechanical behaviour of biocompostable flax/poly(L-lactide) composites, A. Le Duigou, I. Pillin, A. Bourmaud, P. Davies, C. Baley, Composites: Part A, vol. 39, pp 1471–1478 (2008). 80. Polymer-layered silicate nanocomposites: preparation, properties and uses of a new class of materials, M. Alexandre, P. Dubois, Materials Science and Engineering, vol. 28, pp 1–63 (2000). 81. Polymer/layered silicate nanocomposites: a review from preparation to processing, S. S. Ray, M. Okamoto, Prog. Polym. Sci., vol. 28, pp 1539–1641 (2003). 82. An overview on the degradability of polymer nanocomposites, J. K. Pandey, K. R. Reddy, A. P. Kumar, R.P. Singh, Polymer Degradation and Stability, vol. 88, pp 234–250 (2005). 83. Degradation of poly(ethylene terephthalate)/clay nanocomposites during melt extrusion: Effect of clay catalysis and chain extension, X. Xu, Y. Ding, Z. Qian, F. Wang, B. Wen, H. Zhou, S. Zhang, M. Yang Polymer Degradation and Stability, article in press (2009). 84. Morphology and properties of organoclay modified PC/PMMA blend, S. Sinha Ray, M. Bousmina, A. Maazouz, Polymer Engineering and Science, vol. 46, pp 1121–1129 (2006). 85. Thermal degradation behaviour of PE/clay nanocomposites, M. Zanetti, P. Bracco, L. Costa, Polymer Degradation and Stability, vol. 85, pp 657–665 (2004). 86. Photodegradation of PP nanocomposites, B. Mailhot, S. Morlat, J.-L. Gardette, S. Boucard, J. Duchet, J.-F. G´erard, Polymer Degradation and Stability, vol. 82, pp 163–167 (2003). 87. Photooxidation of ethylene-propylene-diene/montmorillonite nanocomposites, S. Morlat-Therias, B. Mailhot, J.-L. Gardette, C. Da Silva, B. Haidar, A. Vidal, Polymer Degradation and Stability, vol. 90, pp 78–85 (2005). 88. Recyclability of a layered silicate-thermoplastic olefin elastomer nanocomposite, M.R. Thomson, K.K. Yeung, Polymer Degradation and Stability, vol. 91, pp 2396–2407 (2006). 89. Effects of reprocessing on the structure of PA6 nanocomposites, I. Goitisolo, J.I. Eguiazabal, J. Nazabal, Polymer Degradation and Stability, vol. 93, pp 1747–1752 (2008). 90. Reprocessing PC/ABS blends: influence on physical properties, J.I. Eguiazabal, J. Nazabal, Polymer Engineering and Science, vol. 30, pp 527–531 (1990).

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91. Recycled PET/layered silicate nanocomposites : morphology and tensile mechanical properties, A. Pegoretti, J. Kolarik, C. Peroni, C. Migliaresi, Polymer, vol. 45, pp 2751–2759 (2004). 92. Facing up to the recycling challenge, G. Marsh, Reinforced plastics (June 2001). 93. On the life cycle metaphor: where ecology and economics diverge, R.U. Ayres, Ecological Economics, vol. 48, n◦ 4, pp 425–438 (2004). 94. FRP recycling in Japan, J. Takahashi, Composites, n◦ 35, pp 58–61 (1999). 95. Development of The Pyrolysis: Process for Recycling of SMC, D.R. Norris, ESD/ASM Advanced Composites conference, Structural Composites: Design and Process Technologies, ASM International, Materials Park, Ohio, pp 277–283 (1990). 96. Pyrolysor diagram, internet site of SITA, http://www.sanifa.fr/pyrolyse.htm (29.10.2009). 97. Recycling of thermoset automotive components, C.N. Cucuras, A.M. Flax, W.D. Graham, G.N. Hartt, SAE Transactions, vol. 100, n◦ 5, pp 437–452 (1991). 98. Recycling automotive related wastes in concrete, G.J. Xu, D.F. Watt, P.P. Hudec, K.A. MacDonald, D.O. Northwood, Journal of Materials Processing Technology, n◦ 48, pp 385–390 (1995). 99. Economics of recycling thermosets, G.N. Hartt, D.P. Carey, SAE Transactions, vol. 101, n◦ 5, pp 636–663 (1992). 100. SMC composites recycling, W.D. Graham, Proceedings of the 9th Annual ASM/ESD Advanced Composites Conference, Detroit MI, USA, pp 609–619 (1993). 101. Thermoset, W.J. Farrissey, Plastics Recycling: Products and Processes, R.J. Ehrig, ed., Hanser Publishers, pp 231–262 (1992). 102. Recycling of sheet-molding compounds by chemical routes, H. Winter, H.A.M. Mostert, G. Paas, J. Appl. Polym. Sci., vol. 57, pp 1409–1417 (1995). 103. Theoretical prediction of weight loss and molecular weight during random chain scission degradation of polymers, K.H. Yoon, I.N. Kang, H.K. Shim, Polymer, vol. 38, pp 2281–2285 (1997). 104. WO Patent 94/25517, T.A. Tufts et al., Ashland Oil, (1994). 105. US Patent 5 468 780, S. Kubota et al., Miyaso Chemical, (1995). 106. EP Patent 0 693 527, S. Kubota et al., Miyaso Chemical, (1996). 107. Studies on the chemical degradation of styrene/vinyl-ester networks, W. Funke et al., Makromol. Chem, vol. 28, p 17 (1958). 108. EP Patent 1 085 044, C.M. Tai et al., Kabushiki Kaisha Toshiba, (2001). 109. Chemical recycling of class A surface quality sheet moulding composites, M. Vall´ee, G. Tersac, N. Destais-Orvoen, G. Durand, Industrial and Engineering Chemistry Research Editorial (2004). 110. SMC composites waste management as reinforcing fillers in polypropylene by combination of mechanical and chemical recycling processes, D. Perrin, C. Guillermain, A. Bergeret, J.M. Lopez-Cuesta, G. Tersac, J. of Materials Science, vol. 41, n◦ 12, pp 3593–3602 (2006). 111. FR Patent 2 728 275 A1, A. Vallet et al., Soci´et´e Europ´eenne de propulsion, (1994). 112. WO Patent 94/26812, A. Piras et al., Cray Valley, (1994). 113. Recycling takes off at MCR, Composites International, n◦ 48, pp 27–31 (2001). 114. Thermosetting materials: recycling and applications, A. Marion, Mecelec Composites & Recyclage, Composites, n◦ 35, pp 49–51 (1999). 115. The implementation of recycled thermoset composites in thermoforming moulds, D.M. Wilson, Journal of Industrial Technology, vol. 19, n◦ 2, pp 45–58 (2003). 116. Final report, A. Molnar, Clean Washington centre (CWC), www.cwc.org (1995). 117. Final report, G. Gaustad, CEER at Alfred University Summer undergraduate Researches Fellowships 2002, Center for Environmental and Energy Research (CEER) (2002). 118. Manufacture of sheets using post-consumer unsaturated polyester resin/glass fibre composites, A.G. Pedroso, D.S. Rosa and T.D.Z., Progress in Rubber, Plastics and Recycling Technology, vol. 18, n◦ 2, pp 111–125 (2002). 119. Recyclage de composites fibers de verre – polyester insatur´e – carbonate de calcium par r´eincorporation dans des matrices thermoplastiques, J.N. Reygrobellet, PhD Dissertation, Universit´e Montpellier II, 234 p. (2000). 120. Treatment of SMC composite waste for recycling as reinforcing fillers in thermoplastics, D. Perrin, E. Leroy, L. Clerc, A. Bergeret, J.M. Lopez-Cuesta, Macromolecular Symposia, vol. 221, pp 227–236, (2005).

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121. Optimizing a recycling process of SMC composite waste, D. Perrin, E. Leroy, L. Clerc, A. Bergeret, J.M. LopezCuesta, Waste Management, 28, pp 541–548 (2008). 122. Recyclage des SMC/BMC – Etude d’un proc´ed´e de revalorization par incorporation des broyats dans une matrice thermoplastique, D. Perrin, PhD Dissertation, Universit´e Montpellier II, 261 p. (2005). 123. Recycled carbon fiber filled polyethylene composites, T Mc Nally, P Boyd, C McClory, D Bien, I Moore, B Millar, J Davidson, T Carroll, Journal of Applied Polymer Science, vol. 107, pp 2015–2021 (2008). 124. N. Rizzi, The waste management of composite materials at the end of the lifetime: technical and economic stakes, Ardi, France.

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25 Nanoparticle Reinforcement of Elastomers and Some Other Types of Polymers James E. Mark Department of Chemistry and the Polymer Research Center, The University of Cincinnati, Cincinnati , Ohio,, USA

25.1 Introduction The main focus of this chapter is characterization of the effects of nanoparticles on the properties of elastomers, but considerable information is provided as well on glassy polymers, partially crystalline polymers, naturally-occurring polymers, relatively rigid polymers, and thermosets. Improvements in mechanical properties are generally of foremost interest, particularly in the case of relatively weak materials, such as poly(dimethylsiloxane). In addition, many fillers such as exfoliated clays also readily provide increased resistance to solvents and decreases in flammability and permeability. Decreases in permeability have become of particular interest with regard to responses to chemical and biological agents. Also, fillers such as expanded graphite and molecular layers of graphene can provide electrical conductivity at remarkably low values of the threshold concentration. Earlier studies in this regard have generally focused on macro composites, such as those containing metallic particles or fibers, or semi-conducting polymers such as polyaniline. Most studies of this type emphasize experiments, but some relevant theory and simulations are described as well. In fact, one of the important unsolved issues in polymer science and engineering is the achievement of a good molecular understanding of how reinforcing particles give the excellent improvements in mechanical properties they often provide. Nanoparticles have been introduced into polymers in a variety of ways. The usual is an ex situ approach in which the particles are generated in a separate first step, and then mechanically blended into a polymer. One of the most interesting newer approaches is an in situ procedure, for example the widely used ‘sol-gel’ approach described in the following section.

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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25.2 Fillers in Elastomers 25.2.1

Generation of Approximately Spherical Particles

A relatively new type of nanocomposite involves the synthesis of ‘ultrastructure’ materials, i.e. materials in which structure can be controlled at the level of around 10 nm. An example of such a synthesis is the ‘sol-gel’ hydrolysis of alkoxysilanes (organosilicates) to give silica, SiO2 [1–3]. The reaction is complicated, involving polymerization and branching, but a typical overall reaction may be written Si(OR)4 + 2H2 O

→

SiO2 + 4ROH

(25.1)

where the Si(OR)4 organometallic species is typically tetraethoxysilane (tetraethylorthosilicate, TEOS). In the ceramics application, the precursor compound is hydrolyzed and then condensed to polymeric chains, the chains become more and more branched, and finally a continuous highly swollen gel is formed. It is first dried at moderately low temperatures to remove volatile species, and then is fired into a porous ceramic object, which can be densified. Not surprisingly, the production of ceramics by this novel route has generated a great deal of interest. The sol–gel hydrolyses and condensations can be carried out within a polymeric matrix to generate nanoparticles of the ceramic material, typically with an average diameter of a few hundred angstroms [1–4]. The polymer typically has end groups, such as hydroxyls, that can participate in the hydrolysis–condensation reactions [5, 6]. Such end groups provide better bonding between the two rather disparate phases, but bonding agents may also be introduced for this purpose. This approach has been used to form ceramic-like phases in a wide variety of elastomers [7–11], as is illustrated in Table 25.1. Important examples are poly(dimethylsiloxane) (PDMS) [Si(CH3 )2 –O–] [7, 8, and hydrocarbon elastomers [9]. The organosilicate that has been most used in this technique is tetraethoxysilane, alias tetraethylorthosilicate (TEOS), with the structure Si(OC2 H5 )4 . The tetramethoxysilane reacts more rapidly than the TEOS, but is seldom used because of toxicity problems. In the case of nonpolar polymers such as the natural rubber and the polyethylene-octene elastomer referenced in Table 25.1, the use of alkyl groups R that are larger than C2 H5 has the advantage of increasing miscibility between the Si(OR)4 and the elastomer. The elastomer that has been studied the most in this regard is PDMS. This is due to this polymer being a relatively weak elastomer in need of reinforcement, particularly with regard to tear strength. It also has the advantage of being capable of easily absorbing the precursor materials generally used in the sol–gel

Table 25.1 Some sol–gel phases precipitated into elastomers. Phases

Elastomers

References

_ Silica, SiO2 “ Titania, TiO2 “ Zirconia, ZrO2 Alumina, Al2 O3 Mixed oxides “

Poly(dimethylsil-oxane) Natural rubber Poly(dimethylsil-oxane) Polyethylene-octene elastomer Poly(dimethylsil-oxane) Poly(dimethylsil-oxane) Poly(dimethylsil-oxane) Polyethylene-octene elastomer

[3, 96] [97, 98] [5] [99] [9, 10, 100] [9, 10, 100] [16] [101]

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process. Considerable reinforcement of elastomers can be achieved in this way. This method for introducing reinforcing particles has a number of advantages over the conventional approach in which separately-prepared filler particles are blended, with difficulty, into the uncrosslinked elastomer before its vulcanization [13, 14]. Because of the nature of the in situ precipitation, the particles are well dispersed and are essentially unagglomerated (as demonstrated by electron microscopy). The mechanism for their growth seems to involve simple homogeneous nucleation and, since the particles are separated by polymer, they do not have the opportunity to coalesce. As referenced in Table 25.1, titania, zirconia, and alumina have also been introduced into PDMS by sol–gel techniques, as have some mixed oxides obtained by co-hydrolysis and co-condensation, for example an organosilicate with an organotitanate. The R groups can be chosen so that one component does not co-hydrolyze and co-condense much more rapidly than the other. Typical transmission electron micrographs of such a filled material show [15] the particles to be relatively ˚ monodisperse, with most of them having diameters in the range 100–200 A. The remainder of this section provides some additional details on the use of this technique [4, 5]. In particular, the methods can be quite general, in that a variety of other precursor materials can be utilized; for example, titanates can be hydrolyzed to titania, aluminates hydrolyzed to alumina, and metal carbonyls photolyzed or thermolyzed to metals or metal oxides. Some ceramic phases other than silica are titania, TiO2 [5], zirconia, ZrO2 [9, 10], alumina, Al2 O3, [9, 10], and mixed oxides [16]. In the sol–gel technique, basic catalysts give precipitated phases that are generally well-defined particles, whereas the acidic catalysts give more poorly-defined, diffuse particles. In some cases, bicontinuous (interpenetrating) phases result. It is also possible to generate the catalyst and other reactants themselves in situ, to give composites of unusually high transparency. It is also possible to polymerize monomers to obtain conducting polymers such as polyaniline within polysiloxane matrices [17]. The method can also be used in a variety of polymers (organic as well as inorganic, non-elastomeric as well as elastomeric). Basic catalysts give precipitated phases that are generally well-defined particles, whereas the acidic catalysts give more poorly-defined, diffuse particles [18]. In some cases, bicontinuous (interpenetrating) phases result [19]. It is also possible to generate the catalyst and other reactants themselves in situ, to give composites of unusually high transparency [20]. Interesting ‘aging’ effects are frequently observed in these systems. If the precipitated particles are left in contact with the hydrolysis catalyst and water they appear to reorganize, so that their surfaces become better defined and their sizes become more uniform [21]. The reinforcing ability of such in situ generated particles has been amply demonstrated for a variety of deformations [4, 5, 9, 22]. In the case of uniaxial extension, the modulus [f * ] frequently increases by more than an order of magnitude, with the isotherms generally showing the upturns at high elongations that are the signature of good reinforcement [23, 24]. Typical results are shown in one portion of Figure 25.1, where α is the extension [25]. As is generally the case in filled elastomers, there is seen to be considerable irreversibility in the isotherms, which is thought to be due to irrecoverable sliding of the chains over the surfaces of the filler particles upon being strained. Analogous results document the reinforcement observed in biaxial extension. These curves exhibit maxima and minima that will be a challenge to those seeking a better molecular understanding of filler reinforcement in general. Some fillers other than silica, for example titania, do give stress-strain isotherms that are reversible, indicating interesting differences in surface chemistry, including increased ability of the chains to slide along the particle surfaces [26]. Such results are also illustrated in Figure 25.1 [26]. Increases in modulus can be lower, but with increase in extensibility. The bonding of PDMS to silica, titania, or silica-titania mixed oxide particles is, however, strong enough to suppress swelling of the polymer, as is illustrated in Figure 25.2 [27]. These results involve equilibrium swelling measurements obtained on unfilled and filled PDMS elastomers to estimate the degree of adhesion between elastomer and filler particles [27]. The results differ greatly from those to be expected for non-adhering fillers, indicating good bonding between the two

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[f *]

Silica

Titania

No filler 0

1

α-1

Figure 25.1 Schematic stress–strain isotherms in elongation for PDMS elastomers unfilled, and reinforced in situ with either silica, titania, or silica-titania mixed oxided.

Vro/Vrf

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φ/(1 – φ) Figure 25.2 Plot of volume fraction ratio Vro/Vrf characterizing the swelling of an unfilled PDMS network relative to that of a filled PDMS network, against filler loading expressed as volume ratio of filler to PDMS φ/(1– φ) (where φ is the volume fraction of filler). The various symbols refer to silica, titania, and silica-titania mixed oxides, but all lie on essentially the same line and all indicate good bonding.

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phases. Obviously, resistance to separation from the surface in such swelling tests does not contradict the chains having considerable mobility along the surfaces of some types of particles, as was just described. A variety of techniques have been used to further characterize these in situ filled elastomers [4, 5]. Density measurements, for example, yield information on the nature of the particles. Specifically, the densities of the ceramic-type particles are significantly less than that of silica itself, and this indicates that the particles presumably contain some unhydrolyzed alkoxy groups or some voids, or both. A number of studies using X-ray and neutron scattering [28, 29] have also been carried out on filled PDMS elastomers [4, 30]. 25.2.2

Glassy Particles Deformable into Ellipsoidal Shapes

It is also possible to obtain reinforcement of a PDMS elastomer by polymerizing a monomer such as styrene to yield hard glassy domains within the elastomer [31]. Roughly spherical polystyrene (PS) particles are formed, as is shown in a portion of Figure 25.3, and good reinforcement is obtained. It is possible to convert the essentially spherical PS particles just described into ellipsoids [5, 32]. First, the PS-PDMS composite is raised to a temperature above the Tg of PS. It is then deformed, and cooled while in the stretched state. The particles are thereby deformed into ellipsoids, and retain this shape when cooled. Uniaxial deformations of the composite give prolate (needle-shaped) ellipsoids, and biaxial deformations give oblate (disc-shaped) ellipsoids [32, 33]. This is shown schematically in Figure 25.3. The prolate ellipsoidal particles represent an interesting intermediate between the roughly spherical particles generally used to reinforce elastomers and the more fiber-like materials used to reinforce thermoplastics. The oblate particles, on the other hand, can be thought of as alternatives to the clay-like platelets exfoliated into polymers to give excellent reinforcements at surprisingly low loadings. Such ellipsoidal particles have been characterized using both scanning and transmission electron microscopy. This gives values for their axial ratios and provides a measure of the extent to which their axes were aligned in the direction of stretching. In these anisotropic materials, elongation moduli in the direction of the

Figure 25.3 The upper sketch represents some originally spherical filler particles being deformed into prolate (needle-shaped) ellipsoids by stretching a polymer matrix in which they reside. This in situ approach also orients the axes of the deformed particles in the direction of the stretching. The lower sketch shows how this orientation can be removed by dissolving away the host matrix. The ellisoidal particles can then be isotropically dispersed within another polymer, as randomized reinforcing particles.

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Prolate, parallel

Spherical

[f *]

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Unfilled 0

α-1

Figure 25.4 Stress–strain isotherms for PDMS-polystyrene composites. The results are for the unfilled PDMS, and the PDMS reinforced with particles that were either spherical or deformed into prolate ellipsoids. In the case of the ellipsoidal particles, one curve corresponds to modulus measurements taken parallel to the direction of the stretching force used to deform (and orient) the particles, an the other was taken perpendicular to this direction.

stretching were found to be significantly larger than those of the untreated PS-PDMS elastomer, whereas in the perpendicular direction they were significantly lower. This is shown in Figure 25.4 [32]. Such differences were to be expected from the anisotropic nature of the systems. In the case of non-spherical particles in general, degrees of orientations are also of considerable importance. One interest here is the anisotropic reinforcements such particles provide, and there have been simulations to better understand the mechanical properties of such composites [34, 35]. The effects of orientation can be removed by dissolving away the host polymer and redispersing the particles isotropically in another elastomer matrix.

25.2.3

Layered Fillers

Exfoliating layered particles such as the clays, mica, or graphite are being used to provide very effective reinforcement of polymers at loading levels much smaller than in the case of solid, roughly spherical particles such as carbon black and silica [36–38]. Some examples are given in Table 25.2. Other properties can also be substantially improved, including increased resistance to solvents, and reduced permeability and flammability [36–42]. Some important examples of elastomers reinforced by clays are PDMS [39] and hydrocarbon elastomers such as natural rubber [37], but there have also been studies of ethylene-propylenediene-monomer elastomers, styrene-butadiene copolymers, styrene-butadiene-styrene block copolymers. acrylonitrile-butadiene rubbers, and acrylonitrile-butadiene-styrene rubbers. The clay most often used is organically-modified montmorillonite, but other types of layered fillers have been used, including ribbon-like laths [43], layered double hydroxides [44], mica [45], and graphite [46]. Additional examples are listed in Table 25.2.

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Table 25.2 Some layered fillers in elastomers. Fillers

Elastomers

References

Montmorillonite “ Layered double hydroxides Bentonite “ Rectorite Fluorohectorite “ Attapulgite Kaolinite-muscoite Graphite “ “

Poly(dimethylsiloxane) Natural rubber EPDM rubber Natural rubber Styrene-butadiene copolymer Styrene-butadiene copolymer Natural rubber, styrene-butadiene copolymer, EPDM rubber Natural rubber Styrene-butadiene copolymer Styrene-butadiene-styrene triblock copolymer Acrylonitrile-butadiene rubber Acrylonitrile-butadiene-styrene rubber EPDM rubber

[39] [37] [102] [103] [104] [105] [106] [107] [108] [109] [110] [111] [112]

25.2.4

Magnetic Particles

Incorporating reinforcing particles that respond to a magnetic field is important with regard to aligning particles that are evenly spherical to improve the mechanical properties of a nanocomposite anisotropically. Some relevant information is given in Table 25.3 [5, 17, 47]. Ferrites and iron oxide particles have been of particular interest in this regard. Considerable anisotropy in structure and mechanical properties can be obtained, as is demonstrated for PDMS [48]. Specifically, the reinforcement is found to be significantly higher in the direction parallel to the magnetic lines of force. Polyaniline has been used to provide some electrical conductivity in some elastomers, and CdSe quantum dots have also been introduced, as referenced in Table 25.3. 25.2.5

Polyhedral Oligomeric Silsesquioxanes (POSS)

These fillers are cage-like silicon-oxygen structures, and have been called the smallest possible silica particles [49]. The most common structure has eight silicon atoms, each carrying an organic group. The particles on which none of the groups are functionally reactive can be simply blended into elastomers such as PDMS using the usual mixing or compounding techniques. POSS molecules having one reactive functional group can be attached to a polymer as side chains. Those with two reactive groups can be incorporated into polymer backbones by copolymerization, and those with more than two can be used for forming crosslinks, yielding network structures. The costs of these particles are gradually decreasing, so they may soon be used more Table 25.3 Electromagnetic phases in elastomers. Types

Examples

References

Magnetic particles “ Electrically semi-conducting polymers Quantum dots

Ferrites γ -Ferric oxide Polyaniline CdSe

[5] [113] [17] [47]

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Handbook of Multiphase Polymer Systems Table 25.4 Polyhedral oligomeric silsesquioxanes (POSS) in elastomers. Bonding functionalities

Polymers

References

1 2 3–8

Poly(dimethylsiloxane) Various (cross-linked) copolymers Various (cross-linked) polymers

[49, 114] [6] [50, 115]

often in both elastomers and non-elastomeric materials. Some relevant studies are described in Table 25.4 [6, 49, 50].

25.2.6

Nanotubes

Carbon nanotubes are also of considerable interest with regard to both reinforcement and possible increases in electrical conductivity [51, 52]. There is considerable interest in characterizing the flexibility of these nanotube structures, in minimizing their tendencies to aggregate, and in maximizing their miscibilities with inorganic as well as organic polymers. Both single-walled and multi-walled nanotubes have been of interest. Some studies of this type are referenced in Table 25.5 [53–55].

25.2.7

Dual Fillers

There can be a considerable advantage to using a combination of fillers of different types, for example particles and layered sheets. A specific example would be carbon black particles and clay layered fillers. One frequently obtains a synergistic effect in that the improvements in properties obtained can be larger than expected from simple additivities. In addition, the first filler may have a solubilizing effect, making incorporation of the second one easier. Relatively little has been done in this regard, but some relevant work is cited in Table 25.6 [56–58].

Table 25.5 Nanotubes in elastomers. Types

Polymers

References

Single-walled Multi-walled “

Styrene-isoprene copolymer Styrene-butyl acrylate copolymer Poly(methylvinylsiloxane)

[116] [117] [118]

Table 25.6 Dual fillers in some nanocomposites. Filler 1 Types

Examples

Filler 2 Types

Examples

Polymers

References

Particles “ “ “

Carbon black Fly ash Carbon black Carbon black

Layered filler “ “ “

Mica Mica Clay Silica

Nitrile rubber Unsaturated polyester Nylon 6 Chlorobutyl rubber

[56] [57] [58] [119]

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Figure 25.5 Polymer chains being threaded through a porous inorganic material such as a zeolite by polymerizing monomer that had been absorbed into one of the channels or cavities.

25.2.8

Porous Fillers

Some fillers are sufficiently porous to accommodate polymer chains, but there is little driving force for a long-chain molecule to reduce its entropy by threading through such an opening. The approach has therefore been to introduce monomers into such cavities, and then polymerizing them. This threads the resulting chains through the cavities, with unusually intimate interactions between the reinforcing phase and the host elastomeric matrix [59–61]. Examples of fillers are some of the zeolites, mesoporous silica, and silica nanotubes, and one of the polymers studied is the elastomer poly(ethyl acrylate). Such an arrangement is illustrated in Figure 25.5, and some relevant references are provided in Table 25.7 [59, 62, 63]. Unusually good reinforcement is generally obtained. Also, because of the constraints imposed by the cavity walls, these confined polymers frequently show no glass transition temperatures or melting points [35]. Similar confinement effects can occur in the case of polymer chains in the galleries of clay particles. 25.2.9

Fillers with Controlled Interfaces

By choosing the appropriate chemical structures, chains that span filler particles in a PDMS-based composite can be designed so that they are either durable, are breakable irreversibly, or are breakable reversibly [64, 65]. Durable chains would be exemplified by those of a polysiloxane, whereas one that would break once at elevated temperatures would contain a peroxide linkage or a linkage of azo-bis-isobutyronitrile. Reversible breakage could be obtained by including hemi-acetal groups somewhere along the chains (see Figure 25.6). 25.2.10

Silicification and Biosilicification

There has been some interest in generating silica-like particles using templates, as is done by Nature in biosilicification processes [66]. Various particle shapes have been obtained, but fibrous forms would be of particular interest if they self assembled into anisotropic structures with anisotropic properties. Platelet forms of silica, on the other hand, would be an interesting synthetic analogue to the naturally-occurring clays, particularly with regard to their possible abilities to provide reinforcement and decreased permeabilities at very low concentrations. Table 25.7 Porous fillers in elastomers. Porous fillers

Polymers

References

Zeolites “ Mesoporous silica Silica nanotubes

Polystyrene, poly(ethyl acrylate) Poly(ethyl acrylate) Polystyrene, poly(ethyl acrylate) Polystyrene, poly(ethyl acrylate)

[59, 62, 120] [63] [59] [59, 121]

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1

R/Si

Figure 25.6 The hardness of a silica-PDMS composite as a function of the relative numbers of alkyl groups and silicon atoms.

25.2.11

Theory and Simulations on Filler Reinforcement

Some relevant theory and simulations are described in Table 25.8 [67–75]. In one approach that has been very useful, the simulations focus on the ways the filler particles change the distribution of the end-to-end vectors of the polymer chains making up the elastomeric network, from the fact that the filler nanopartices exclude the chains from the volumes they occupy [70]. The changes in the polymer chain distributions from this filler ‘excluded volume effect’ then cause associated changes in the mechanical properties of the elastomer host matrix. Specifically, Monte Carlo simulations were carried out as has been done for the unfilled networks, but now each bond of the chain was tested for overlapping with a filler particle as the chain was being generated [76]. If any bond penetrated a particle surface, the entire chain conformation was rejected and a new chain started. The particle sizes of greatest interest are those used commercially, with small particles giving significantly better reinforcement than larger ones. The primary particles are generally assumed to be spherical. In actual filled elastomers, the particles are dispersed at least relatively randomly, but it was of interest to do simulations on regular particle arrangements as well [76]. Table 25.8 Some relevant theory and simulations on nanocomposites. Approach

Goal

References

Analytical theory

Reinforcement

[67–69]

Monte Carlo simulations

Chain conformations around filler particles, and reinforcement

[70, 71, 80]

Molecular dynamics simulations

“

[72–74]

“

Chain dynamics

[122, 123]

Self-consistent field theory with Markov chain statistics

Dispersion of polymers into clay galleries

[75]

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In one such study [77], the reinforcing particles were randomly distributed. Of greatest interest in these studies was whether the particles cause increases or decreases in the end-to-end distances, with this expected to depend particularly on the sizes of the filler particles, but presumably on other variables such as their concentration in the elastomeric matrix as well [77]. One effect of the particles was to increase the dimensions of the chains when the filler particles were small relative to the dimensions of the network chains. In contrast, particles that were relatively large tended to decrease the chain dimensions. These simulated results on the distributions are in agreement with some subsequent neutron scattering experiments on deuterated and non-deuterated chains of PDMS [78]. The polymers contained silica particles that were surface treated to make them inert to the polymer chains, as was implicitly assumed in the simulations. These experimental results also indicated chain extensions when the particles were relatively small, and chain compressions when they were relatively large. In a rather different approach, Mattice et al. [79] generated particles within a matrix by collapsing some of the chains into domains that would act as reinforcing filler. They found that small particles did lead to significant increases in chain dimensions, while large particles led to moderate decreases, in agreement with the single-chain simulations and scattering experiments. Additional simulations eliminated the possibility that the collapse of some of the chains gave rise to an increase in free volume into which the non-collapsed chains expanded [80]. The cases where the filler causes compression of the chain are relevant to another area of rubberlike elasticity, specifically the preparation of networks by crosslinking in solution followed by removal of the solvent [5]. Such experiments were initially carried out to obtain elastomers that had fewer entanglements and the success of this approach was supported by the observation that such networks came to elastic equilibrium much more rapidly. They also exhibited stress–strain isotherms in elongation that were closer in form to those expected from the simplest molecular theories of rubberlike elasticity. The most recent work in this area has focused on the unusually high extensibilities of such elastomers [81–83]. In any case, the present simulations should help elucidate molecular aspects of phenomena in this research area as well. There are two items of primary interest here, specifically increases in modulus in general, and upturns in the modulus with increasing deformation. Results are typically expressed as the reduced nominal or engineering stress as a function of deformation. The area under such curves up to the rupture point of the sample then gives the energy of rupture, which is the standard measure of the toughness of a material [5]. The stress–strain isotherms in elongation [77] from such simulated distributions showed substantial increases in modulus that increased with increase in filler loading, as expected. Additional increases would be expected by taking into account other mechanisms for reinforcement such as physisorption, which could be modeled by Lennard-Jones interactions between the particles and the chains. Treating the case of chemisorption would be more complicated, but one could model randomly-distributed active particle sites that were interacting very strongly with the chains by means of a Dirac δ-function. When the distance between a chain and an active site became less than the range of short-range interactions, then the chain would be considered chemisorbed. Simulations have also been carried out on ellipsoidal particles such as the polystyrene prolate and oblate ellipsoids mentioned above. For example, oriented prolate particles [84] showed that the anisotropy in structure causes the values of the modulus in the longitudinal direction to be significantly higher than those in the transverse direction. These simulated results are in at least qualitative agreement with the experimental differences in longitudinal and transverse moduli obtained experimentally [32]. In spite of their inherent interest, relatively little has been done on oblate ellipsoidal fillers. In one study, however, oblate particles were again placed on a cubic lattice [34], and were oriented in a way consistent with their orientation in PS-PDMS composites that were the subject of an experimental investigation [33]. In general, the network chains tended to adopt more compressed configurations relative to those of prolate

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a

b

c

Figure 25.7 Sketches of (a) primary particles, (b) aggregates, and (c) agglomerates occurring in fillers such as carbon black and silica.

particles having equivalent sizes and aspect ratios. The elongation moduli were found to depend on the sizes, number, and axial ratios of the particles, as expected. In particular, the reinforcement from the oblate particles was found to be greatest in the plane of the particles, and the changes were in at least qualitative agreement with the corresponding experimental results [33]. It would obviously be of interest to model various distributions of sizes, shapes, and orientations. The silica or carbon black particles used to reinforce commercial materials are seldom completely dispersed [13, 85–88], as is assumed in the simulations described. As is shown schematically in Figure 25.7, the primary particles are generally aggregated into relatively stable aggregates and these aggregates are frequently clustered into less stable arrangements called agglomerates. Simulations are being carried out on such more highly ordered structures [89]. The shapes of the aggregates being included are linear, globular, branched, star-shaped, and fractal. It is well known in the industry that such structures are important in maximizing the reinforcement, as evidenced by the fact that being too persistent in removing such aggregates and agglomerates in blending procedures gives materials with less than optimal mechanical properties [13, 85–88]. Friedlander et al. have demonstrated that such aggregates have a remarkable deformability, by carrying out elongation experiments both reversibly, and irreversibly to their rupture points [90–93]. This is of considerable importance, since when these structures are within elastomeric matrices, their deformations upon macroscopic deformations of the filled elastomer means that they must contribute to the storage of the elastic deformation energy. This would have to be taken into account both in the interpretation of experimental results and in more refined simulations of filler reinforcement.

25.3 Nanoparticles in Glassy Polymers This type of thermoplastic maintains rigidity when the temperature is below the glass transistion temperature Tg . Particularly important examples of polymers in this category are polystyrene, polycarbonate, and poly(methyl methacrylate). These and other glassy polymers have been reinforced by a variety of nanoparticles, including silica, titania, POSS, and some of the clays. As is the case with non-glassy polymers, the nanocomposites with graphite are of particular interest because of the electrical conductivity they frequently exhibit at very low values of the percolation threshold concentration. Some typical studies of nanocomposites based on polymers of this type are referenced in Table 25.9.

25.4 Nanoparticles in Partially-Crystalline Polymers This second class of thermoplastic polymers maintains rigidity from crystallites present when the temperature is below the melting temperature Tm . The polyamides or Nylons, polyethylene, and isotactic polypropylene

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Table 25.9 Some results on nanoparticles in glassy thermoplastics. Nanoparticles

Polymers

References

Silica “ Titania Various clays Montmorillonite Nanotubes “ “ “ Graphite “ “ “ Polyhedral oligomeric silsesquioxane (POSS)

Polystyrene Unsaturated polyester Various Polystyrene Foamed polysulfone Polycarbonate Poly(methyl methacrylate) Polystyrene Foamed polystyrene Poly(methyl methacrylate) Polystyrene Unsaturated polyester Polycarbonate Polynorbornene

[9, 10] [124] [125] [36] [42] [55] [53, 54, 126] [54, 127] [128] [129, 130] [129, 131] [132] [133] [134, 135]

have been of particular interest in this regard. The nanoparticles have included silica, various clays, nanotubes, graphite, and POSS. Information on representative studies of such polymers of this type is presented in Table 25.10. How well dispersed nanoparticles are in these, and other, polymers is a matter of some contention [94].

25.5 Nanoparticles in Naturally-Occurring Polymers Polymers in this category are of particular current interest since they are generally also biodegradable. They include proteins, chitosan, and poly(3-hydroxyalkanoates) such as copolymers of poly(3-hydroxybutyrate), Table 25.11 describes some studies on members in this important category. Environmental concerns worldwide are likely to make this an increasingly important class of polymers [95], particularly in nanocomposites. Their preparations, characterizations, and applications are clearly in the category of ‘green’ chemistry. Table 25.10 Some results on nanoparticles in partially crystalline thermoplastics. Nanoparticles

Polymers

References

Silica Titania Clays Layered double hydroxides Nanotubes “ “ Graphite “ Polyhedral oligomeric silsesquioxane (POSS)

Nylons Various Nylons Polyethylene Polyethylene Polypropylene Poly(vinyl alcohol) Polyethylene Polypropylene Polyethylene

[9, 136] [125] [40] [137] [54, 127, 138] [139, 140] [54] [46, 141–143] [144] [135, 145]

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Table 25.11 Some results on nanoparticles in naturally-occurring polymers. Nanoparticles

Polymers

References

Silica “ Titania “ Zirconia Montmorillonite Polyhedral oligomeric silsesquioxane (POSS) “

Linseed oil Proteins Drying oils Hydroxypropylcellulose Drying oils Chitosan Poly(3-hydroxyalkanoates) Proteins

[11, 146] [147] [148, 149] [150] [149] [41] [151] [152]

25.6 Nanoparticles in Relatively-Rigid Polymers These polymers are already high-performance materials, so the levels of reinforcement obtainable can be significantly lower than those shown by weaker polymers. In other words, it’s hard to improve materials that already have very impressive properties. Some information on such polymers is given in Table 25.12. The polyimides represent the most extensively studied polymers in this category, and the sol–gel approach has been used to reinforce them and other rigid polymers, at least to some extent. In some cases, it is difficult to find convenient solvents for dissolution of this type of polymer. Another important problem is getting good interfacial bonding between the nanoparticles and polymers that are designed in part to be unreactive. Because of their structures, many of these polymers can be electrically conductive, at least in the doped state. Some may also exhibit nonlinear optical properties.

25.7 Nanoparticles in Thermoset Polymers Thermosets are heavily crosslinked polymers that have solidity and good mechanical properties by virtue of all the covalent bonds locking in a permanent network structure. The epoxies have been the most studied, but various thermosets based on formaldehyde have also been of interest. Some studies on nanoparticles in such polymers are described in Table 25.13. Table 25.12 Some results on nanoparticles in relatively-rigid polymers. Nanoparticles

Polymers

References

Silica “ “ Titania Alumina Montmorillonite Nanotubes Polyhedral oligomeric silsesquioxane (POSS) “

Poly(amide imides) Polyimides Polyimides, benzobisoxazole polymers Polyimides Polyimides Triphenylphosphate polymer Poly(p-phenylene benzobisoxazole) Polyimides Liquid-crystalline materials

[153] [154] [10] [155] [156] [38] [157] [50, 158] [115]

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Table 25.13 Some nanoparticles in thermosets. Nanoparticles

Polymers

References

Silica Titania Montmorillonite “ “ “

Epoxy Epoxy Epoxy Urea-formaldehyde resin Melamine-formaldehyde resin Polyimide Polycyanurate Epoxy Epoxy Epoxy Epoxy Epoxy Polyimide Epoxy

[159, 160] [161] [162] [163] [164] [165] [166] [167] [168] [169] [170] [161, 171] [172] [135]

Layered double hydroxide Mica Sepiolite Nanotubes Graphite “ Polyhedral oligomeric silsesquioxane (POSS)

This network structure makes these polymers very resistant to aggressive environments, but also makes them non-reprocessible.

25.8 Conclusions The documentation cited clearly demonstrates the successful use of nanoparticles to reinforce elastomers and some other types of polymers. The stage is certainly set for others to enter this interesting and challenging area of polymer science and engineering, and to contribute by further investigating the properties and applications of some of these interesting and important nanocomposites.

25.9 Acknowledgements It is a pleasure to acknowledge the financial support provided by the National Science Foundation through Grants DMR-0314760 DMR-0803454 (Polymers Program, Division of Materials Research).

References 1. A.I. Nakatani, R.J. Hjelm, M. Gerspacher, and R. Krishnamoorti (Eds), Filled and Nanocomposite Polymer Materials, Materials Research Society, Warrendale, PA, 2001, Vol. 661. 2. Sanchez, R.M. Laine, S. Yang, and C.J. Brinker (Eds), Organic/Inorganic Hybrid Materials – 2002, C. Materials Research Society, Warrendale, PA, 2002), vol. 726. 3. J.E. Mark, Some Novel Polymeric Nanocomposites, Acct. Chem. Res., 39, 881–888 (2006). 4. J.E. Mark and D.W. Schaefer, Reinforcement of Elastomers by the In-Situ Generation of Filler Particles, in PolymerBased Molecular Composites, D.W. Schaefer and J.E. Mark (Eds), Materials Research Society, Pittsburgh, 1990, Vol. 171, pp. 51–56.

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5. B. Erman and J.E. Mark, Structures and Properties of Rubberlike Networks. Oxford University Press, New York, 1997. 6. J.E. Mark, Some Unusual Elastomers and Experiments on Rubberlike Elasticity, Prog. Polym. Sci., 28, 1205–1221 (2003). 7. J.E. Mark, Some Interesting Things About Polysiloxanes, Acct. Chem. Res., 37, 946–953 (2004). 8. J.E. Mark, H.R. Allcock, and R. West, Inorganic Polymers, 2nd ed. Oxford University Press, New York, 2005. 9. J.E. Mark, The sol-gel route to inorganic-organic composites, Hetero. Chem. Rev., 3, 307–326 (1996). 10. J.E. Mark, Ceramic-Reinforced Polymers and Polymer-Modified Ceramics, Polym. Eng. Sci., 36, 2905–2920 (1996). 11. Z. Zong, M.D. Soucek, and C. Xue, Unusual Inorganic Phase Formation in Ultraviolet-Curable Organic-Inorganic Hybrid Films, J. Polym Sci. A, Polym. Chem., 43, 1607–1623 (2005). 12. D. Fragiadakis and P. Pissis, Glass Transition and Segmental Dynamics in Poly(dimethylsiloxane)/Silica Nanocomposites Studied by Various Techniques, J. Non-Cryst. Solids, 353, 4344–4352 (2007). 13. Z. Rigbi, Reinforement of Rubber by Carbon Black, Adv. Polym. Sci., 36, 21–68 (1980). 14. A.I. Medalia and G. Kraus, Reinforcement of Elastomers by Particulate Fillers, in Science and Technology of Rubber, J.E. Mark, B. Erman, and F.R. Eirich (Eds), Academic, San Diego, 1994, pp. 387–418. 15. Y.-P. Ning, M.-Y. Tang, C.-Y. Jiang, J.E. Mark, and W.C. Roth, Particle Sizes of Reinforcing Silica Precipitated Into Elastomeric Networks, J. Appl. Polym. Sci., 29, 3209–3212 (1984). 16. J. Wen and J.E. Mark, Synthesis, Structure, and Properties of Poly(dimethylsiloxane) Networks Reinforced by in Situ-Precipitated Silica-Titania, Silica-Zirconia, and Silica-Alumina Mixed Oxides, J. Appl. Polym. Sci., 58, 1135–1145 (1995). 17. D. Zhou, S. Subramaniam, and J.E. Mark, In-Situ Synthesis of Polyaniline in Poly(dimethylsiloxane) Networks Using an Inverse Emulsion Route, J. Macro. Sci.-Chem., 42, 113–126 (2005). 18. K.D. Keefer, Structure and Growth of Silica Condensation Polymers, in Silicon-Based Polymer Science, in SiliconBased Polymer Science. A Comprehensive Resource, J.M. Zeigler and F.W.G. Fearon (Eds), Am. Chem. Soc., Washington, DC, 1990, Vol. 224, pp. 227–240. 19. D.W. Schaefer, L. Jian, C.-C. Sun, D.W. McCarthy, C.-Y. Jiang, Y.-P. Ning, J.E. Mark, and S. Spooner, StructureProperty Relationships in Silica-Siloxane Molecular Composites, in Ultrastructure Processing of Advanced Materials, D.R. Uhlmann and D.R. Ulrich (Eds), Wiley, New York, 1992, pp. 361–375. 20. G.S. Rajan, G.S. Sur, J.E. Mark, D.W. Schaefer, and G. Beaucage, Preparation and Characterization of Some Unusually Transparent Poly(dimethylsiloxane) Nanocomposites, J. Polym Sci. B, Polym. Phys., 41, 1897–1901 (2003). 21. P. Xu, S. Wang, and J.E. Mark, Particle Growth and Development During the In-Situ Precipitation of Silica in a Polymeric Matrix, in Better Ceramics Through Chemistry IV, B.J.J. Zelinski, C.J. Brinker, D.E. Clark, and D.R. Ulrich (Eds), Materials Research Society, Pittsburgh, 1990, Vol. 180, pp. 445–451. 22. J.E. Mark and B. Erman, Rubberlike Elasticity. A Molecular Primer. Wiley-Interscience, New York, 1988. 23. J.E. Mark and Y.-P. Ning, Effects of Ethylamine Catalyst Concentration in the Precipitation of Reinforcing Silica Filler in an Elastomeric Network, Polym. Bulletin, 12, 413–417 (1984). 24. J.E. Mark, S. Wang, P. Xu, and J. Wen, Reinforcement from In-Situ Precipitated Silica in Polysiloxane Elastomers Under Various Types of Deformation, in Submicron Multiphase Materials, R.H. Baney, L.R. Gilliom, S.-I. Hirano, and H.K. Schmidt (Eds), Materials Research Society, Pittsburgh, PA, 1992, Vol. 274, pp. 77–84. 25. S. Wang, P. Xu, and J.E. Mark, Shear and Biaxial Extension Measurements of Reinforcement from In-Situ Precipitated Silica, Rubber Chem. Technol., 64, 746–759 (1991). 26. S.-B. Wang and J.E. Mark, In-Situ Precipitation of Reinforcing Titania Fillers, Polym. Bulletin, 17, 271–277 (1987). 27. J. Wen and J.E. Mark, Precipitation of Silica-Titania Mixed-Oxide Fillers into Poly(Dimethylsiloxane) Networks, Rubber Chem. Technol., 67, 806–819 (1994). 28. R.-J. Roe, Methods of X-Ray and Neutron Scattering in Polymer Science. Oxford University Press, Oxford, 2000. 29. G.D. Wignall, Small-Angle-Neutron-Scattering Characterization of Polymers, in Physical Properties of Polymers. Third Edition, J.E. Mark, K.L. Ngai, W.W. Graessley, L. Mandelkern, E.T. Samulski, J.L. Koenig, and G.D. Wignall (Eds), Cambridge University Press, Cambridge, 2004, pp. 424–511.

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30. D.W. Schaefer, J.E. Mark, D.W. McCarthy, L. Jian, C.-C. Sun, and B. Farago, Structure of Microphase-Separated Silica/Siloxane Molecular Composites, in Polymer-Based Molecular Composites, D.W. Schaefer and J.E. Mark (Eds), Materials Research Society, Pittsburgh, 1990, Vol. 171, pp. 57–63. 31. F.-S. Fu and J.E. Mark, Polystyrene-Polyisobutylene Network Composites from In-Situ Polymerizations, J. Appl. Polym. Sci., 37, 2757–2768 (1989). 32. S. Wang and J.E. Mark, Generation of Glassy Ellipsoidal Particles within an Elastomer by In-situ Polymerization, Elongation at an Elevated Temperature, and Finally Cooling Under Strain, Macromolecules, 23, 4288–4291 (1990). 33. S. Wang, P. Xu, and J.E. Mark, Method for Generating Oriented Oblate Ellipsoidal Particles in an Elastomer and Characterization of the Reinforcement They Provide, Macromolecules, 24, 6037–6039 (1991). 34. M.A. Sharaf and J.E. Mark, Monte Carlo Simulations on Filler-Induced Network Chain Deformations and Elastomer Reinforcement from Oriented Oblate Particles, Polymer, 43, 643–652 (2002). 35. J.E. Mark, New Developments and Directions in the Area of Elastomers and Rubberlike Elasticity, Macromol. Symp., Kyoto issue, 201, 77–83 (2003). 36. T.J. Pinnavaia and G. Beall (Eds), Polymer-Clay Nanocomposites, Wiley, New York, 2001. 37. Y.T. Vu, J.E. Mark, L.H. Pham, and M. Engelhardt, Clay Nanolayer Reinforcement of Cis-1,4-Polyisoprene and Epoxidized Natural Rubber, J. Appl. Polym. Sci., 82, 1391–1403 (2001). 38. W. Zhou, J.E. Mark, M.R. Unroe, and F.E. Arnold, Some Clay Nanocomposites Based on a High-Temperature, High-Performance Polymer, J. Macromol. Sci. – Pure Appl. Chem., A38, 1–9 (2001). 39. D.F. Schmidt, F. Clement, and E.P. Giannelis, On the Origins of Silicate Dispersion in Polysiloxane/Layered-Silicate Nanocomposites, Adv. Functional Mater., 16, 417–425 (2006). 40. A. Usuki, N. Hasegawa, and M. Kato, Polymer-Clay Nanocomposites, Adv. Polym. Sci., 179, 135–195 (2005). 41. Y. Xu, X. Ren, and M.A. Hanna, Chitosan/Clay Nanocomposite Film Preparation and Characterization, J. Appl. Polym. Sci., 99, 1684–1691 (2006). 42. H. Sun, J.E. Mark, S.C. Tan, N. Venkatasubramanian, M.D. Houtz, F.E. Arnold, and C.Y.-C. Lee, Microcellular Foams from Some High-Performance Thermoplastics and Their Composites, Nonlinear Opt., Quantum Opt., 31, 1–29 (2004). 43. T.-M. Wu, Y.-H. Lien, and S.-F. Hsu, Isothermal Crystallization Kinetics and Melting Behavior of Nylon/Saponite and Nylon/Montmorillonite Nanocomposites, J. Appl. Polym. Sci., 94, 2196–2204 (2004). 44. F.R. Costa, U. Wagenknecht, D. Jehnichen, M. Abdel-Goad, and G. Heinrich, Nanocomposites Based on Polyethylene and Mg-Al Layered Double Hydroxide. II. Rheological Characterization, Polymer, 47, 1649–1660 (2006). 45. S. Bose and P.A. Mahanwar, Influence of Particle Size Distribution on Mica Filled Nylon 6 Composite, J. Mats. Sci., 40, 6423–6428 (2005). 46. W. Zheng, X. Lu, and S.-C. Wong, Electrical and Mechanical Properties of Expanded Graphite-Reinforced HighDensity Polyethylene, J. Appl. Polym. Sci., 91, 2781–2788 (2004). 47. G.A. Ozzin and A.C. Arsenault, Nanochemistry. A Chemical Approach to Nanomaterials. RSC Publlishing, London, 2005. 48. G.B. Sohoni and J.E. Mark, Anisotropic Reinforcement in Elastomers Containing Magnetic Filler Particles, J. Appl. Polym. Sci., 34, 2853–2859 (1987). 49. G. Pan, J.E. Mark, and D.W. Schaefer, Synthesis and Characterization of Fillers of Controlled Structure Based on Polyhedral Oligomeric Silsesquioxane Cages, and Their Use in Reinforcing Siloxane Elastomers, J. Polym. Sci., Polym Phys Edn., 41, 3314–3323 (2003). 50. S.H. Phillips, T.S. Haddad, and S.J. Tomczak, Developments in Nanoscience: Polyhedral Oligomeric Silsesquioxanes (POSS)-Polymers, Curr. Opinion Solid State Mater. Sci., 8, 21–29 (2004). 51. J.D. Lichtenhan, J. Schwab, and W.A. Reinerth, Sr., Nanostructured Chemicals: A New Era in Chemical Technology, Chem. Innov., 31, 3–5 (2001). 52. R. Tamaki, J. Choi, and R.M. Laine, A Polyimide Nanocomposite from Octa(aminophenyl)silsesquioxane, Chem. Mater., 15, 793–797 (2003). 53. F. Du, J.E. Fischer, and K.I. Winey, Coagulation Method for Preparing Single-Walled Carbon Nanotube/Poly(methyl methacrylate) Composites and Their Modulus, Electrical Conductivity, and Thermal Stability, J. Polym. Sci., Polym. Phys. Ed., 41, 3333–3338 (2003).

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54. R. Andrews and M.C. Weisenberger, Carbon Nanotube Polymer Composites, Curr. Opinion Solid State Mater. Sci., 8, 31–37 (2004). 55. P. Potschke, H. Brunig, A. Janke, D. Fischer, and D. Jehnichen, Orientation of Multi-Walled Carbon Nanotubes in Composites with Polycarbonate by Melt Spinning, Polymer, 46, 10355–10363 (2005). 56. N. Nugay and B. Erman, Property Optimization in Nitrile Rubber Composites via Hybrid Filler Systems, J. Appl. Polym. Sci., 79, 366–374 (2001). 57. S. Sen and N. Nugay, Tuning of Final Performance of Unsaturated Polyester Composites with Inorganic Microsphere/Platelet Hybrid Reinforcers, Eur. Polym J., 37, 2047–2053 (2001). 58. Y. Konishi and M. Cakmak, Structural Heirarchy Developed in Injection Molding of Nylon 6/Clay/Carbon Black Nanocomposites, Polymer, 46, 4811–4826 (2005). 59. H.L. Frisch and J.E. Mark, Nanocomposites Prepared by Threading Polymer Chains through Zeolites, Mesoporous Silica, or Silica Nanotubes, Chem. Mater., 8, 1735–1738 (1996). 60. J.E. Mark, Polymer Chains in Constraining Environments, in Molecular Catenanes, Rotaxanes and Knots, J.-P. Sauvage and C. Dietrich-Buchecker (Eds), Wiley-VCH, Weinheim, 1999, pp. 223–246. 61. G. Kickelbick, Concepts for the Incorporation of Inorganic Building Blocks into Organic Polymers on a Nanoscale, Prog. Polym. Sci., 28, 83–114 (2003). 62. H.L. Frisch, Y. Xue, S. Maaref, G. Beaucage, Z. Pu, and J.E. Mark, Pseudo Interpenetrating Polymer Networks and Interpenetrating Polymer Networks of Zeolite 13X and Polystyrene and Polyethylacrylate, Macromol. Symp., 106, 147–166 (1996). 63. H.L. Frisch, S. Maaref, Y. Xue, G. Beaucage, Z. Pu, and J.E. Mark, Interpenetrating and Pseudo-Interpenetrating Polymer Networks of Polyethylacrylate and Zeolite 13X, J. Polym. Sci., Polym. Chem. Ed., 34, 673–677 (1996). 64. D.W. Schaefer, B.T.N. Vu, and J.E. Mark, The Effect of Interphase Coupling on the Structure and Mechanical Properties of Silica-Siloxane Composites, Rubber Chem. Technol., 75, 795–810 (2002). 65. B.T.N. Vu, J.E. Mark, and D.W. Schaefer, Interfacial Modification for Controlling Silica-Polysiloxane Interactions and Bonding in Some Elastomeric Composites, Composite Interfaces, 10, 451–473 (2003). 66. S.V. Patwardhan, V.P. Taori, M. Hassan, N.R. Agashe, J.E. Franklin, G. Beaucage, J.E. Mark, and S.J. Clarson, An Investigation of the Properties of Poly(dimethylsiloxane)-Bioinspired Silica Hybrids, Eur. Polym J., 42, 167–178 (2006). 67. T.A. Witten, M. Rubinstein, and R.H. Colby, Reinforcement of Rubber by Fractal Aggregates, J. Phys. II France, 3, 367–383 (1993). 68. G. Heinrich, M. Kluppel, and T.A. Vilgis, Reinforcement of Elastomers, Curr. Opinion Solid State Mats. Sci., 6, 195–203 (2002). 69. M. Kluppel, The Role of Disorder in Filler Reinforcement of Elastomers on Various Length Scales, Adv. Polym. Sci., 164, 1–86 (2003). 70. J.E. Mark, R. Abou-Hussein, T.Z. Sen, and A. Kloczkowski, Some Simulations on Filler Reinforcement of Elastomers, Polymer, 46, 8894–8904 (2005). 71. M. Vacatello, Predicting the Molecular Arrangements in Polymer-Based Nanocomposites, Macromol. Theor. Sims., 12, 86–91 (2003). 72. J.B. Hooper, K.S. Schweizer, T.G. Desai, R. Koshy, and P. Keblinski, Structure, Surface Excess and Effective Interactions in Polymer Nanocomposite Melts and Concentrated Solutions, J. Chem. Phys., 121, 6986–6997 (2004). 73. D. Gersappe, Molecular Mechanisms of Failure in Polymer Nanocomposites, Phys. Rev. Lett., 89, 58301–5830158304 (2002). 74. T. Desai, P. Keblinski, and S.K. Kumar, Computer Simulations of the Conformations of Strongly Adsorbed Chains at the Solid-Liquid Surface, Polymer, 47, 722–727 (2006). 75. A.C. Balazs, C. Singh, E. Zhulina, and Y. Lyatskaya, Modeling the Phase Behavior of Polymer/Clay Nanocomposites, Acc. Chem. Res., 32, 651–657 (1999). 76. J.E. Mark, Some Simulations on Filler Reinforcement in Elastomers, Molec. Cryst. Liq. Cryst., 374, 29–38 (2002). 77. Q.W. Yuan, A. Kloczkowski, J.E. Mark, and M.A. Sharaf, Simulations on the Reinforcement of Poly(dimethylsiloxane) Elastomers by Randomly-Distributed Filler Particles, J. Polym. Sci., Polym. Phys. Ed., 34, 1647–1657 (1996).

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78. A.I. Nakatani, W. Chen, R.G. Schmidt, G.V. Gordon, and C.C. Han, Chain Dimensions in Polysilicate-Filled Poly(Dimethylsiloxane), Polymer, 42, 3713–3722 (2001). 79. F.M. Erguney, H. Lin, and W.L. Mattice, Dimensions of Matrix Chains in Polymers Filled with Energetically Neutral Nanoparticles, Polymer, 47, 3689–3695 (2006). 80. F.M. Erguney and W.L. Mattice, Response of Matrix Chains tol Nanoscale Filler Particles, Polymer, 49, 2621–2623 (2008). 81. K. Urayama and S. Kohjiya, Uniaxial Elongation of Deswollen Polydimethylsiloxane Networks with Supercoiled Structure, Polymer, 38, 955–962 (1997). 82. S. Kohjiya, K. Urayama, and Y. Ikeda, Poly(Siloxane) Network of Ultra-High Elongation, Kautschuk Gummi Kunststoffe, 50, 868–872 (1997). 83. K. Urayama and S. Kohjiya, Extensive Stretch of Polysiloxane Network Chains with Random- and Super-Coiled Conformations, Eur. Phys. J. B, 2, 75–78 (1998). 84. M.A. Sharaf, A. Kloczkowski, and J.E. Mark, Monte Carlo Simulations on Reinforcement of an Elastomer by Oriented Prolate Particles, Comput. Theor. Polym. Sci., 11, 251–262 (2001). 85. B.B. Boonstra, Role of Particulate Filers in Elastomer Reinforcement: A Review, Polymer, 20, 691–704 (1979). 86. E.L. Warrick, O.R. Pierce, K.E. Polmanteer, and J.C. Saam, Silicone Elastomer Developments 1967-1977, Rubber Chem. Technol., 52, 437–525 (1979). 87. J.-B. Donnet and A. Vidal, Carbon Black: Surface Properties and Interactions with Elastomers, Adv. Polym. Sci., 76, 103–127 (1986). 88. J. Donnet and E. Custodero, Reinforcement of Elastomers by Particulate Fillers, in Science and Technology of Rubber, J.E. Mark and B. Erman (Eds), Elsevier, Amsterdam, 2005, pp. 367–400. 89. R. Abou-Hussein and J.E. Mark, unpublished results. 90. S.K. Friedlander, H.D. Jang, and K.H. Ryu, Elastic Behavior of Nanoparticle Chain Aggregates, Appl. Phys. Lett., 72, 11796–11799 (1998). 91. Y.J. Suh, S.V. Prikhodko, and S.K. Friedlander, Nanostructure Manipulation Device for Transmission Electron Microscopy: Application to Titania Nanoparticle Chain Aggregates, Microsc. Microanal., 8, 497–501 (2002). 92. Y.J. Suh and S.K. Friedlander, Origin of the Elastic Behavior of Nanoparticle Chain Aggregates: Measurements Using Nanostructure Manipulation Device, J. Appl. Phys., 93, 3515–3523 (2003). 93. W. Rong, A.E. Pelling, A. Ryan, J.K. Gimzewski, and S.K. Friedlander, Complementary TEM and AFM Force Spectroscopy to Characterize the Nanomechanical Properties of Nanoparticle Chain Aggregates, Nano Lett., 4, 2287–2292 (2004). 94. D.W. Schaefer and R.S. Justice, How Nano are Nanocomposites? Macromolecules, 40, 8502–8517 (2007). 95. A.L. Andrady (Ed), Plastics and the Environment, Wiley-Interscience, New York, 2003. 96. J.E. Mark and B. Erman, Rubberlike Elasticity. A Molecular Primer, 2nd Edition. Cambridge University Press, Cambridge, 2007. 97. S. Kohjiya, A. Katoh, J. Shimanuki, T. Hasegawa, and Y. Ikeda, Three-Dimensional Nano-Structure of In Situ Silica in Natural Rubber as Revealed by 3D-TEM/electron Tomography, Polymer, 46, 4440–4436 (2005). 98. Y. Ikeda, S. Poompradub, Y. Morita, and S. Kohjiya, Preparation of HIgh Performance Nanocomposite Elastomer: Effect of Reaction Conditions on In Situ Silica Content in Natural Rubber, J. Sol-Gel Sci. Techn., 45, 299–306 (2008). 99. H.-T. Liao and C.-S. Wu, Organic-Inorganic Polymeric Nanocomposites Involving Titanium Tetraisopropylate inf Polyethylene-Octene Elastomer, J. Polym. Sci., Polym. Phys., 42, 4272–4280 (2004). 100. N. Yamada, I. Yoshinaga, and S. Katayama, Formatiion Behavior and Optical Proerties of Transparent InrganicOrganic Hybrids Prepared from Metal Alkoxides and Polydimethylsiloxane, J. Sol-Gel Sci. Techn., 17, 123–130 (2000). 101. C.-S. Wu, Synthesis of Polyethylene-Octene Elastomer/SiO2 -TiO2 Nanocomposites via In Situ Polymerization: Properties and Characterization of the Hybrid, J. Polym. Sci., Polym. Chem., 43, 1690–1701 (2005). 102. H. Acharya, S.K. Srivastava, and A.K. Bhowmick, Synthesis of Partially Exfoliated EPDM/LDH Nanocomposites by Solution Intercalation: Structural Characterization and Properties, Comp. Sci. Tech., 67, 2807–2816 (2007). 103. M.A. Lopez-Manchado, M. Arroyo, B. Herrero, and J. Biagiotti, Vulcanization Kinetics of Natural RubberOrganoclay Nanocomposites, J. Appl. Polym. Sci., 89, 1–15 (2003).

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104. Y. Wang, L. Zhang, C. Tang, and D. Yu, Preparation and Characterization of Rubber-Clay Nanocomposite, J. Appl. Polym. Sci., 78, 1879–1883 (2000). 105. Y. Wang, H. Zhang, Y. Wu, J. Yang, and L. Zhang, Preparation, Structure, and Properties of a Novel Rectorite/StyreneButadiene Copolymer Nanocomposite, J. Appl. Polym. Sci., 95, 324–328 (2005). 106. Y.-P. Wu, Y. Ma, Y.-Q. Wang, and L.-Q. Zhang, Effects of Characteristics of Rubber, Mixing and Vulcanization on the Structure and Properties of Rubber/Clay Nanocomposites by Melt Blending, Macromol. Mater. Eng., 289, 890–894 (2004). 107. S. Varghese and J. Karger–Kocsis, Natural Rubber-Based Nanocomposites by Latex Compounding with Layered Silicates, Polymer, 44, 4921–4927 (2003). 108. M. Tian, L. Cheng, W. Liang, and L. Zhang, The Anisotropy of Fibrillar Silicate/Rubber Nanocomposites, Macromol. Mater. Eng., 290, 681–687 (2005). 109. Z. Chen and K. Gong, Preparation and Dynamic Mechanical Proerties of Poly(styrene-b-butadiene)-Modified Clay Nanocomposite, J. Appl. Polym. Sci., 84, 1499–1503 (2002). 110. J. Yang, M. Tian, Q.-X. Jia, L.-Q. Zhang, and X.-l. Li, Influence of Graphite Particle Size and Shape on the Properties of NBR, J. Appl. Polym. Sci., 102, 4007–4015 (2006). 111. H.S. Dahiya, N. kishore, and R.M. Mehta, Effect of Percolation on Electrical and Dielectric Proerties of Acrylonitrile Butadiene Styrene/Graphite Composite, J. Appl. Polym. Sci., 106, 2101–2110 (2007). 112. E. Planes, J. Duchet, A. Naazouz, and J.-F. Gerard, Characterization of New Formulations for the Rotational Molding Based on Ethylene-Propylene Copolymer/Graphite Nanocomposites, Polym. Eng. Sci., 48, 489–498 (2008). 113. N. Guskos, V. Likodimos, S. Glenis, M. Maryniak, M. Baran, R. Szymczak, Z. Roslaniec, M. Kwiatkowska, and D. Petridis, Magnetic Properties of γ - Fe2 O3 /Poly(ether-ester) Nanocomposites, J. Nanosci. Nanotech., 8, 2127–2134 (2008). 114. A. Striolo, C. McCabe, and P.T. Cummings, Thermodynamic ad Transport Properties of Polyhedral Oligomeric Silsesquioxanes in Poly(dimethylsiloxane), J. Phys. Chem., B, 109, 14300–14307 (2005). 115. R.M. Laine, Nanobuilding Blocks Based on the [OSiO1.5 ]x (x = 6, 8, 10) Octasilsesquioxanes, J. Mats. Chem., 15, 3725–3744 (2005). 116. H.J. Baraza, F. Pompeo, E.A. O’Rear, and D.E. Resasco, SWNT-Filled Thermoplastic and Elastomeric Composites Prepared by Miniemulsion Polymerization, Nano Lett., 2, 795–802 (2002). 117. A. Dufresne, M. Paillet, J.L. Putaux, R. Canet, F. Carmona, P. Delhaes, and S. Cui, Processing and Characterization of Carbon Nanotube/Poly(styrene-co-butyl acrylate) Nanocomposites, J. Mats. Sci., 37, 3915–3923 (2002). 118. M.-J. Jiang, Z.-M. Dang, and H.-P. Xu, Enhanced Electrical Conductivity in Chemically Modified Carbon Nanotube/Methylvinyl Silicone Rubber Nanocomposite, Eur. Polym J., 43, 4924–4930 (2007). 119. V. Sridhar, R.N.P. Chaudhary, and D.K. Tripathy, Relaxation Behaior of Carbon Silica Dual Phase Filler Reinforced Chlorobutyl Vulcanizates, J. Appl. Polym. Sci., 101, 4320–4327 (2006). 120. S. Maaref, H.L. Frisch, G.S. Rajan, Z. Pu, J.E. Mark, and G. Beaucage, Glassy Polystyrene Composites Having Zeolite 13X or Vycor as a Dispersed Inorganic Phase, J. Macromol. Sci., Pure Appl. Chem., A36, 1895 (1999). 121. H.L. Frisch, J.M. West, C.G. Goltner, and G.S. Attard, PS in nanotubes, J. Polym. Sci., Polym. Chem. Ed., 34, 1823 (1996). 122. D. Brown, V. Marcadon, P. Mele, and N.D. Alberola, Effect of Particle Size on the Properties of Model Nanocomposites, Macromolecules, 41, 1499–1511 (2008). 123. T.G. Desai, P. Keblinski, and S.K. Kumar, Polymer Chain Dynamics at Interfaces: Role of Boundary Conditions at Solid Interface, J. Chem. Phys., 41, 044903–044901-044907 (2008). 124. C.-M. Chung, S.-Y. Cho, J.-G. Kim, and S.Y. Oh, Preparation of Unsaturated Polyester-Silica Nanocomposites, J. Appl. Polym. Sci., 106, 2442–2447 (2007). 125. B. Dunn and J.I. Zink (Eds), Special issue, 40(9), on Sol-Gel Chemistry and Materials, Acct. Chem. Res., (2007). 126. R.E. Gorga and R.E. Cohen, Toughness Enhancements in Poly(methyl methacrylate) byAddition of Oriented Multiwall Carbon Nanotubes, J. Polym. Sci., Polym. Phys., 42, 2690–2702 (2004). 127. M.C. Weisenberger, R. Andrews, and T. Rantell, Carbon Nanotube Polymer Composites: Recent Developments in Mechanical Properties, in Physical Properties of Polymers Handbook, J.E. Mark (Eds), Springer-Verlag, New York, 2006, pp. 585–598.

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128. J. Shen, C. Zeng, and L.J. Lee, Synthesis of Polystyrene-Carbon Nanofibers Nanocomposite Foams, Polymer, 46, 5218–5224 (2005). 129. G.-H. Chen, D.-J. Wu, W.-G. Weng, and W.-L. Yan, Preparation of Polymer/Graphite Conducting Nanocomposite by Intercalation Polymerization, J. Appl. Polym. Sci., 82, 2506–2513 (2001). 130. W.-P. Wang, Y. Liu, X.-X. Li, and Y.-Z. You, Synthesis and Characteristics of Poly(methyl methacrylate)/Expanded Graphite Nanocomposites, J. Appl. Polym. Sci., 100, 1427–1431 (2006). 131. G.-H. Chen, D.-J. Wu, W.-G. Weng, B. He, and W.-L. Yan, Preparation of Polystyrene-Graphite Conducting Nanocomposites via Intercalation Polymerization, Polym. Int., 50, 980–985 (2001). 132. W. Lu, H. Lin, D. Wu, and G. Chen, Unsaturated Polyester Resin/Graphite Nanosheeet Conducdting Composites with a Low Percolation Threshold, Polymer, 47, 4440–4444 (2006). 133. M. Abdel-Goad, P. Potschke, D. Zhou, J.E. Mark, and G. Heinrich, Preparation and Rheological Characterization of Polymer Nanocomposites Based on Expanded Grapite, J. Macromol. Sci., Pure Appl. Chem., 44, 591–598 (2007). 134. H.G. Jeon, P.T. Mather, and T.S. Haddad, Shape Memory and Nanostructure in Poly(norbornyl-POSS) Copolymers, Polym. Int., 49, 453–457 (2000). 135. G. Pan, Polyhedral Oligomeric Silsesquioxane (POSS), in Physical Properties of Polymers Handbook, J.E. Mark (Ed), Springer-Verlag, New York, 2006, pp. 577–584. 136. S. Jain, H. Goosens, M. van Duin, and P. Lemstra, Effect of In Situ Prepared Silica Nano-Particles on Non-Isothermal Crystallization of Polypropylene, Polymer, 46, 8805–8818 (2005). 137. F.R. Costa, M. Saphiannikova, U. Wagenknecht, and G. Heinrich, Layered Double Hydroxide Based Polymer Nanocomposites, Adv Polym. Sci., 210, 101–168 (2008). 138. S.L. Ruan, P. Gao, X.G. Yang, and T.X. Yu, Toughening High Performance Ultrahigh Molecular weight Polyetahylene Using Multiwalled Nanotubes, Polymer, 44, 5643–5654 (2003). 139. A.R. Bhattacharyya, T.V. Sreekumar, T. Liu, S. Kumar, L.M. Ericson, R.H. Hauge, and R.E. Smalley, Crystallization and Orientation Studies in Polypropylene/Single Wall Carbon Nanotube Composite, Polymer, 44, 2373–2377 (2003). 140. E. Assouline, A. Lustiger, A.H. Barber, C.A. Cooper, E. Klein, E. Wachtel, and H.D. Wagner, Nucleation Ability of Multiwall Carbon Nanotubes in Polypropylene Composites, J. Polym. Sci., Polym. Phys. Edn., 41, 520–527 (2003). 141. M. Alexandre, M. Pluta, P. Dubois, and R. Jerome, Metallocene Catalyzed Polymerization of Ethylene in the Presence of Graphite, 1, Macromol. Chem. Phys., 202, 2239–2246 (2001). 142. G. Wu, J. Lin, Q. Zheng, and M. Zhang, Correlation between Percolation Behavior of Electricity and Viscoelasticity for Graphite Filled High Density Polyethylene, Polymer, 47, 2442–2447 (2006). 143. J. Lu, X. Chen, W. Lu, and G. Chen, The Piezoresistive Behavior of Polyethylene/Foliated Graphite Nanocomposites, Eur. Polym J., 42, 1015–1021 (2006). 144. V. Causin, C. Marega, A. Marigo, G. Ferrara, and A. Ferraro, Morphological and Structural Characterization of Polypropylene/Conductive Graphite Nanocommposites, Eur. Polym J., 42, 3153–3161 (2006). 145. K. Pielichowski, J. Njuguna, B. Janowski, and J. Pielichowski, Polyhedral Oligomeric Silsesquioxnaes (POSS) Containing Nanohybrid Polymers, Adv. Polym. Sci., 201, 225–296 (2006). 146. M.D. Soucek, Inorganic/Organic Hybrid Coatings, in Hybrid Materials. Synthesis, Characterization, and Applications, G. Kickelbick (Ed), Wiley-VCH Verlag, Weinheim, 2007, pp. 433–476. 147. T. Coradin and J. Livage, Aqueous Silicates in Biologicl Sol-Gel Applications: New Perspectives for Old Precursors, Acct. Chem. Res., 40, 819–826 (2002). 148. C.R. Wold, H. Ni, and M.D. Soucek, Model Reaction Study on the Interaction between the Inorganic and Organic Phases in Drying Oil Based Ceramer Coatings, Chem. Mater., 13, 3032–3037 (2001). 149. R.L. Ballard, J.P. Williams, J.M. Njus, B.R. Kiland, and M.D. Soucek, Inorganic-Organic Hybrid Coatings with Mixed Metal Oxides, Eur. Polym J., 37, 381–398 (2001). 150. M. Kusabe, H. Kozuka, S. Abe, and H. Suzuki, Sol-Gel Preparation and Properties of HydroxypropylcelluloseTitania Hybrid Thin Films, J. Sol-Gel Sci. Technol., 44, 111–118 (2007). 151. R. Hany, R. Hartmann, C. Bohlen, S. Brandenberger, J. Kawada, C. Lowe, M. Zinn, B. Witholt, and R.H. Marchessault, Chemical Synthesis and Characterization of POSS-Functionalized Poly(3-hydroxyalkanoates), Polymer, 46, 5025–5031 (2005).

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152. R.Y. Kannan, H.J. Salacinski, P. E.Butler, and A.M. Seifalian, Polyhedral Oligomeric Silsesquioxane Nanocomposites: The Next Generation Material for Biomedical Applications, Acc. Chem. Res., 38, 879–884 (2005). 153. K. Babooram, B. Franiis, R. Bissessur, and R. Narain, Synthesis and Characterization of Novel (Amide-Imide)-Silica Composites by the Sol-Gel Process, Comp. Sci. Tech., 68, 617–624 (2008). 154. M. Khalil, S. Saeed, and Z. Ahmad, Mechanical and Thermal Properties of Polyimide/Silica Hybrids with ImideModified Silica Network Structure, J. Appl. Polym. Sci., 107, 1257–1268 (2008). 155. S. Koytepe, T. Seckin, N. Kivrilcim, and H.I. Adiguzel, Synthesis and Dielectric Properties of Polwyimide-Titania Hybrid Composites, J. Inorg. Organomet. Polym. Mater., 18, 222–228 (2008). 156. P. Ma, W. Nie, Z. Yang, P. Zhang, G. Li, Q. Lei, L. Gao, X. Ji, and M. Ding, Preparation and Characterization of Polyimide/Al2O3 Hybrid Films by Sol-Gel Process, J. Appl. Polym. Sci., 108, 705–712 (2008). 157. S. Kumar, T.D. Dang, F.E. Arnold, A.R. Bhattacharyya, B.G. Min, X. Zhang, R.A. Vaia, C. Park, W.W. Adams, R.H. Hauge, R.E. Smalley, S. Ramesh, and P.A. Willis, Synthesis, Structure, and Properties of PBO/SWNT Composites, Macromolecules, 35, 9039–9043ww (2002). 158. Y.-S. Ye, W-Y.Chen, and Y.-Z. Wang, Synthesis and Properties of Low-Dielectric-Constant Polyimides with Introduced Reactive Fluorine Polyhedral Oligomeric Silsesquioxanes, J. Polym Sci. A, Polym. Chem., 44, 5391–5402 (2006). 159. W. Araki and T. Adachi, Viscoelastsicity of Epoxy Resin/Silica Hybrid Material Prepared via Sol-Gel Process: Considered in Terms of Morphology, J. Appl. Polym. Sci., 107, 253–261 (2007). 160. W. Araki, S. Wada, and T. Adachi, Viscoelastsicity of Epoxy Resin/Silica Hybrid Materials with an Acid Anhydride Curing Agenta, J. Appl. Polym. Sci., 108, 2421–2427 (2008). 161. G. Xian, R. Walter, and F. Haupert, A Synergistic Effect of Nano-TiO2 and Graphite on the Tribological Performance of Epoxy Matrix Composites, J. Appl. Polym. Sci., 102, 2391–2400 (2006). 162. V. Mittal, Effect of the Presence of Excess Amonium Ions on the Clay Surface on Permeation Properties of Epoxy Nanocomposites, J. Mater. Sci., 43, 4972–4978 (2008). 163. H. Lei, G. Du, A. Pizzi, and A. Celzard, Influence of Nanoclay on Urea-Formaldehyde Resins for Wood Adhesives and Its Model, J. Appl. Polym. Sci., 109, 2442–2451 (2008). 164. H. Wang, X. Meng, Z. Qian, H. Zhou, Y. Ding, S. Zhang, and M. Yang, Synthesis and Properties of MelamineFormaldehyde/Montmorillonite Nanocomposites, J. Nanosci. Nanotech., 8 1775–1781 (2008). 165. M.J. Ginert, S.C. Jana, and S.G. Miller, An Optimum Treatment of Nanoclay for PMR-15 Nanocomposites, Polymer, 48, 7573–7581 (2007). 166. G.I. Anathoulis, E. Kontou, A. Fainleib, I. Bei, and Y. Gomza, Synthesis and Characterization of Polycyanurate/Montmorillonite Nanocomposites, J. Polym. Sci., Polym. Phys., 46, 1036–1049 (2008). 167. C.-H. Tseng, H.-B. Hsueh, and C.-Y. Chen, Effect of Reactive Layered Double Hydroxides on the Thermal and Mechanical Properties of LDHs/Epoxy Nanocomposites, Comp. Sci. Tech., 67, 2350–2362 (2007). 168. K. Tamura, S. Yokoyama, C.S. Pascua, and H. Yamada, New Age of Polymer Nanocomposites Containing Dispersed High-Aspect-Ratio Silicate Nanolayers, Chem. Mater., 20, 2242–2246 (2008). 169. A. Nohales, L. Solar, I. Porcar, C.I. Vallo, and C.M. Gomez, Morphology, Flexural, and Thermal Properties of Sepiolite Modified Epoxy Resins with Different Curing Agents, Eur. Polym. J., 42, 3093–3101 (2006). 170. J.K.W. Sandler, J.E. Kirk, I.A. Kinloch, M.S.P. Shaffer, and A.H. Windle, Ultra-Low Electrical Percolation Threshold in Carbon-Nanotube-Epoxy Composites, Polymer, 44, 5893–5899 (2003). 171. J. li, M.L. Sham, J.-K. Kim, and G. Marom, Morphology and Properties of UV/OzoneTreated Graphite Nanoplatelet/Epoxy Nanocomposites, Comp. Sci. Tech., 67, 293–305 (2007). 172. D. Cho, S. Lee, G. Yang, H. Fukushima, and L.T. Drzal, Dynamic Mechanical and Thermal Properties of Phenylethynyl-Terminated Polyimide Composites Reinforced with Expanded Graphite Nanoplatelets, Macromol. Mater. Eng., 290, 170–187 (2005).

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Index

ABC miktoarm star copolymers 38 abrasion characteristics 9 absorption cross sections 709 acetolysis 928 acid-base interactions 615–16 acorn morphology 194–5 acoustic emission 782–3 acrylamide (AAm) 895–6 2-acrylamido-2-methyl propane sulfonic acid (AMPS) 275–6 acrylated epoxidized soybean oil (AESO) 911 acrylic acid (AA) 208 acrylonitrile-butadiene rubber (NBR) 125, 253–5, 269 acrylonitrile-butadiene-styrol (ABS) 377, 565 activating interlayers 101 active films 764–6 active-PAMPS gels 149–50 Adam-Gibbs theory 480–1, 492 additive flame retardancy 845, 849–62 additive migration 803–6, 826, 829 adhesion 84–5, 87, 768, 961–3 adhesives 458, 609–12, 886–7, 891 adsorption–desorption processes 84, 750 see also gas transport mechanisms aerogels 145 aeronautics applications 882–91 AFM see atomic force microscopy agar microgels 144 Agari model 413–14 ageing processes 10, 797–841 additive migration 803–6, 826, 829 application-related factors 798, 799, 815–16, 832–9 associated issues 797–8 chemical ageing 798, 806–29 classification 798–9 definitions 797–8 diffusion controlled kinetics 805–6, 831–2 elastomers 961 end-of-life criteria 798, 799 environmental factors 797, 798–9, 836–7, 839

evaporation controlled kinetics 804–5 mechanistic schemes 816–29 monitoring systems 886 multiphase polymer systems 829–32, 837–9 physical ageing 798, 799–806, 832–9 solvent absorption 802–3 structural reorganization 800–2, 829–31 superficial layer 815–16 agglomerates 433–5, 437, 672–3 aggregates 433–5, 453, 455–6 aircraft structural health monitoring (ASHM) 886 alloys 251 alumina 853–4, 859, 961 aluminum fillers 436 aluminum nitride 407, 414 aluminum tri-hydroxide (ATH) 849–50, 859–61, 881, 932 amine crosslinked epoxy (ACE) 823 aminolysis 946–7 3-aminopropyltriethoxysilane 145 ammonium persulfate (APS) 144–5 ammonium polyphosphate (APP) 379, 851, 855, 857, 859 amorphous polymer phases ageing processes 800–2, 829–31, 838 characterization 522–6, 530–1, 536, 546, 640, 660–1, 669–70, 673–8 interlayers 90–2, 100–1, 111–12 thermophysical properties 395, 401 amphiphilic gel-type resins 565 analytical theories 49–55 annealing processes 453 Ansatz model 36 antimony trioxide 854–5, 927 antioxidants 803, 806, 820, 824–6, 933–4 applied thermal excitation 786–7 artificial photosynthesis 104–5 aspect ratios 2, 430, 445, 755, 757, 760 atom transfer radical polymerization (ATRP) 99, 592, 600–1, 604–5, 624–6 atomic force microscopy (AFM) 8, 109–10, 142, 202–3, 656

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas. © 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.

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atomistic models 7–8, 58 ATRP see atom transfer radical polymerization attapulgites (ATT) 650–1 auto shredder residues (ASR) 931 automated tape layer (ATL) process 887 automotive applications 866–71 avalanche breakdown 834 back small angle light scattering (BSALS) 656–8 backscattered neutrons 721–3 ballistic protection 879–80 barium titanate 403–4, 406–7 barrier multiphase materials 749, 752–67 barrier properties see gas transport mechanisms batch mixers 128–9, 166–7 Bauer model 289–90 Beer-Lambert equation 592 BEM see boundary element method Benveniste–Miloh model 408 BGY lattice theory 42 BIM see boundary integral method binding energy (BE) 587–91 binodal stability 40 bioactive polymer composites 897, 899, 902, 904–5 bioadhesive drug carriers 544 bioavailability 81 biochemical ageing 797 biocompatibilization 93–100 biocomposites 375, 939–40 biodegradability 764, 895, 898–900, 902–3, 939–40 bioelectronics 104–5 biomimetic surface engineering 101–3 biosilicification 967 bipolarons 451 bismaleimide (BMI) 628 bladder molding 875–6 blob models 36–8 block copolymer-epoxy (BC/E) blends 280 block copolymers applications 900 capillary extrusion rheometry 347–53 dynamic viscoelasticity 340–5 elastomers 964–5, 966 electron spin resonance spectroscopy 561–9 flow alignment 345–7 flow-induced morphological changes 345–7 interfacial properties 99, 110 manufacturing techniques 149 mechanical properties 261–3, 279–81

microphase-separated 312, 339–40, 343 morphological characteristics 176–7, 198–9, 203–4 neutron scattering 730, 731–3 rheological properties 312–13, 339–54 self-consistent field theory 45–52, 68–9 solid-state NMR spectroscopy 526–30, 542–3, 545–6 X-ray photoelectron spectroscopy 628–9 blowing agents 852 bond percolation 427 bone prosthesis applications 895–905 Born approximation 708–10, 713, 716, 734–5 boron-reinforced epoxy composites 875, 882 boundary element method (BEM) 171 boundary integral method (BIM) 171 bovine serum albumin (BSA) 603–4 bow-tie probes 376 Bragg equation/law 677–8, 707–8, 722 Bragg peaks 674–5, 728, 733 Bragg reflections 712–13, 717, 728, 734, 737–9 Bragg scattering 718, 721 bridge decks 878–9 broadband dielectric spectroscopy 483–5, 487, 493–4, 496, 503–7 brominated bisphenols 843, 850 Brownian dynamics 63 Bruggeman model 405 BSALS see back small angle light scattering bulk molding compound (BMC) 871, 924, 948–9 bulk properties 9–10, 65–6 butterfly effect 824 butterfly scattering patterns 731 butylmethacrylate (BMA) 155 cadmium selenide quantum dots 965 Cahn-Hilliard linear theory 647, 650 calcium carbonate fillers 922, 923, 931–2, 937 calorimetry 392–3 capillary extrusion rheometry 347–53 capillary instabilities 172, 319, 320–1 carbohydrates 530 carbon black (CB) elastomers 966, 970 electrical properties 430, 432–5, 453–4, 457 morphological characteristics 226 responsive interphases 104 waste management 922, 932 carbon black filled polyethylene (CBPE) 799 carbon black–poly(vinyl pyrrolidone) (CB–PVP) 912 carbon dioxide, supercritical 132–3 carbon/epoxy composites 867, 872–6, 888–9 carbon fiber reinforced plastics (CFRP) 295–6

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Index carbon fiber reinforcement applications 867–9, 872–9, 885, 892, 898–9, 903 waste management 923, 929, 951 carbon fiber sheet molding compound (CFSMC) 868 carbon fiber/carbon matrix (C/C) composites 889 carbon–glass fiber-reinforced plastic (CFRP) 873–4, 878–9, 882, 885 carbon/glass hybrids 877–8 carbon molecular sieves 767–9 carbon nanofiber nanocomposites (CNF) 142 carbon nanoparticles (CN) 143 carbon nanotubes (CNT) applications 892–3, 908, 910 characterization 507–12, 561–2 elastomers 966 electrical properties 443–9 endohedral filling 449 fire retardancy 853, 859 interfacial properties 94 mechanical properties 265, 300–1 oxidative treatments 447–8 surface modifications 446–9 carbonization 836–7, 851–2, 859 2-carboxyethyl(phenylphosphinic) acid (HPPPA) 861 carboxyl-terminated butadiene acrylonitrile (CTBN) copolymer 279–80 carrier layers 101–2 catalytic conversion 931 cellulose acetate (CA) 275–6 cellulose fibers 451 ceramic matrix composites (CMC) 778 ceramics 879, 899, 903–4, 908, 960–1 cetylpyridinium chloride (CPC) 574 CF see linear correlation function chain correlations 42 chain dynamics 38–9 chain scission 808–12, 823–4 chain welding 808–9, 812–13, 823–5 chaotic mixing 221–7 Charlesby’s theory 809 chemical ageing 798, 806–29 application-related factors 815–16 diffusion processes 813–15 hydrolytic ageing 826–9 intrinsic chemical stability 817–18 lateral group changes 806–8 macromolecular skeleton changes 808–13 mechanistic schemes 816–29 oxidation reactions 806–8, 812, 814–15, 818–24, 839 superficial layer 815–16 chemical compatibilization 93–100, 102

983

chemical gels 659 chemical interfacial interactions 85 chemical recycling 928, 946–7 chirality 444–5 chitosan 144, 763, 897–8 clay nanocomposites 4 electron spin resonance spectroscopy 575–6 fire retardancy 852–3, 857–9, 861 future developments 11 gas transport mechanisms 761–3 interfacial properties 94–5, 102 light scattering techniques 660 mechanical properties 291–3, 299–300, 302–3 polymer gels 146, 148 waste management 934, 940–1, 943 see also polymer–clay nanocomposites click chemistry 99–100 closed loop stabilization 824–5 cloud-point curves 651–2 CMC see ceramic matrix composites; critical micelle concentration co-continuous morphologies 208–11 light scattering techniques 650 mechanical properties 253–5, 270–2, 274 rheological properties 334, 336–9 coarse-grained models 7–8, 33–4, 56–61, 70 coefficient of thermal expansion (CTE) 9, 911 coherent scattering function 711, 712 coincidence detection ERD 111 cold crystallization 829–30 Cole–Cole diagrams 491–2, 494 colloidal structures 605–8, 714 combustibility 846–9 combustion of polymers 844–5, 846–9 combustion of waste materials 929–32 compatibilization compatibilizer effects 263–5 compatibilizer roles 67–8, 198–201 droplet breakup and deformation 201–5 electrical properties 455–6 fire retardancy 857, 860–2 gas transport mechanisms 763 interfacial properties 87, 93–100, 102 mechanical properties 260–5 microfibrillar reinforced composites 671, 681–93 morphological characteristics 197–208, 213–14 polymer blend nanocomposites 207–8, 213–14 processing effects 205 ternary and quaternary blends 205–7 waste management 943 complete encapsulation 193

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complex electric moduli 484–5 complex flows 221 Complex Langevin method 61–2 complex permittivity 426 compliance 93 composition correlations 42 compression molding 453 computed tomography (CT) 786 conducting polymer-coated fillers 451 conducting polymers see inherently conducting polymers conductive grafts 600 conductive heat transfer 388 conductive rubbers 458 cone calorimetry 848–9, 857, 860 confined copolymer melts 69 conformational analysis 522–5 construction applications 866–82 automotive applications 866–71 ballistic protection 879–80 corrosion repair 877 energy industry applications 865, 877–8 infrastructure 878–9 marine applications 871–5 mass transit applications 880–2 pressure vessels 876 sporting goods 875–6 contact angle measurement 142 contact resistance 469–72 continuous chaotic advection blenders (CCAB) 223 continuous mixers 129 continuous path models 36 continuous wave (CW) ESR technique 559–60, 567 continuum Gaussian chains 35 continuum modeling 7–8 continuum viscoelasticity 13 contrast variation 730 controlled interface fillers 967 controlled interfacial morphology 90–2 controlled radical polymerization (CRP) 600 convection heat transfer 388 converging channels 219 converting interphases 102–3 cooperativity length 480 copolyamide (CPA) 323 copolymers applications 900 conformations 53–5 elastomers 964–5, 966 electrical properties 448, 456 electron spin resonance spectroscopy 560–9 light scattering techniques 647, 661

manufacturing techniques 124–33, 139, 145, 149, 154 mechanical properties 254–6, 261–3, 275, 279–81 modeling and simulations 45–52, 67–70 morphological characteristics 176–7, 198–9, 203–4 neutron scattering 730, 731–3 rheological properties 312–13, 339–45 solid-state NMR spectroscopy 526–30, 542–3, 545–6 structure–property relationships 479–80, 482, 486–7, 493–9 waste management 934–5 X-ray photoelectron spectroscopy 589–90, 618–19, 628–9 copper fillers 436 copper ion determinations 601–2 Coran–Patel model 268, 275–6 core–shell structures interfacial properties 98–9 morphological characteristics 195–7, 210–11 thermophysical properties 408–10 X-ray photoelectron spectroscopy 605–6 correction functions 463–6, 471 corrosion repair 877 cost effectiveness 123 Couette flow cells 168, 317 Cox-Merz rule 323–4 CP see cross-polarization CP-2EPOX additive 128–9 crack formation 812, 815–16, 832–5, 890 crazing 802 creep curves 800–2 critical micelle concentration (CMC) 574 critical volume fraction (CVF) 429 cross-polarization (CP) 519–20, 521–2, 525–36, 537 crosslinked proton-conducting membranes (CPM) 531 crosslinking ageing processes 808–9, 812–13, 823–5 characterization 376–7, 571, 731 elastomers 961 interfacial properties 87, 102 interpenetrating polymer networks 154–5, 282–3, 286–7 mechanical properties 260, 272–3, 277–8, 282–3, 286–7, 290–1 morphological characteristics 201 polymer blends 260, 272–3, 277–8 polymer gels 143–9, 290–1 waste management 927, 935 CRP see controlled radical polymerization cryogenic micronization 935 cryogenic properties 901 crystal size 676, 677–8

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Index crystallinity ageing processes 797, 800–2, 829–31, 838 electrical properties 454–5 gas transport mechanisms 756 light scattering techniques 640, 660–2 medical applications 905 nanofillers 970–1 neutron scattering 726–9, 733–43 solid-state NMR spectroscopy 522–6, 530–1, 536, 540, 543, 546 thermophysical properties 395, 401 waste management 943 X-ray scattering 669–70, 673–8, 695 crystallizable polyurethane (CPU) 154 crystallization ageing processes 800–2, 829–31 characterization 8–9, 360–9, 373, 640, 660–72, 733–43 interfacial properties 87, 90 isothermal 361–3 manufacturing techniques 124 mechanical properties 259, 262–3, 266 modeling and simulations 38–9 near surface 733–43 rheological properties 322 Cussler–Lape model 759, 760–1 Cussler’s model 759–60 CW see continuous wave cyclam-functionalized poly(glycidyl methacrylate) chains (CF-PGMA-Cy) 601–2 cyclic loading generated thermal excitation 787 Davies model 268, 271, 275–6 DDFT see dynamic density functional theories de Gennes’ theory 15–16, 41, 51–2 deca-bromodiphenyl ether (deca-BDE) 927 DEER see pulsed double electron–electron resonance deformability see compliance degradation see ageing processes degree of crystallinity 674–6 dehydrochlorination 934 dehydroxylation 934 delamination 780, 782, 786, 793, 888, 890 dense melts 38 dense polymer films 749–51 density functional theory (DFT) 47, 64–5 dental composites 865, 905–7 depolymerization reactions 807, 818 DFT see density functional theory diaryl derivatives 554 dibutyl phthalate (DBP) 429, 434

985

4-[(5-dichloromethylsilyl)pentyloxy] cyanobenzene (DCN) 537 DICO method 394 dicumyl peroxide (DCP) 260–1, 272–5, 277–8 dielectric breakdown 834, 837–8 dielectric properties 10, 425–6 dielectric relaxation spectroscopy (DRS) 483–5, 487, 493–4, 496, 503–7 dielectric spectroscopy 479–517 broadband techniques 483–5, 487, 493–4, 496, 503–7 carbon nanotubes 507–12 copolymers 479–80, 482, 486–7, 493–9 dynamic mechanical analysis 480, 487–9, 492–4, 500–1, 504–7, 513 electrically conductive polymer nanocomposites 507–12 glass transition temperature 480–1, 490, 492, 497 interfacial phenomena 499–507 interpenetrating polymer networks 479–80, 482, 486–95 mixing and phase separation 486–99 molecular dynamics 480, 482, 499–500, 513 morphological characteristics 480, 482, 499 percolation phenomena 507–12 polymer blends 479–80 polymer composites/nanocomposites 479–80, 482, 499–513 responsive interphases 108–9 rubber/silica nanocomposites 499–507 structure–property relationships 479–80 techniques 482–6 theoretical background 482–3 thermally stimulated depolarization current analysis 481–2, 485–9, 493–507, 512–13 diethanolamine (DEA) 610 2-(diethylamino)ethyl methacrylate (DEA) 629 diethylenetriamine (DETA) 946–7 differential scanning calorimetry (DSC) 9, 360, 365–77 ageing processes 801, 825, 829–30 glass transition temperature 481, 558 manufacturing techniques 133, 145, 151, 153–4 molecular dynamics 480, 499–500, 513 rheological properties 345 temperature modulated 373–7 thermophysical properties 392–3 differential scattering cross sections 709–11 diffuse reflectance 922 diffusion ageing processes 802–6, 813–15, 823, 826, 831–2 gas transport mechanisms 749, 750, 754, 766

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diffusion (Continued) interfacial properties 92 rheological models 24–5 diffusion controlled kinetics 805–6, 831–2 diffusivity 9, 10, 389, 390–2, 393–4 diglycidyl ether of bisphenol A (DGEBA) 303, 609–12, 626 dipalmitoylphosphatidylcholine (DPPC) 578 1,1-diphenyl-2-picrylhydrazyl (DPPH) radicals 567 dipolar decoupling 521–2, 525, 536–7 direct mixing method (DMM) 300–1 discrete Gaussian chain model 34–5 discrimination induced by variable angle minipulse (DIVAM) 537–8 dispersion mechanisms 163–4, 178–82, 200, 752–63, 769 dispersive mixing 452 dissipative particle dynamics (DPD) 60, 63–4 dissolution processes 314–15, 924–7 distributive mixing 452–3 DIVAM see discrimination induced by variable angle minipulse divinyl benzene (DVB) 146 DLS see dynamic light scattering DMA see dynamic mechanical analysis DMTA see dynamic mechanical thermal analysis DNA adsorption isotherms 616–18 Doi-Edwards reptation model 24–5 doping 450–1 double-diaphragm forming 887–8 double emulsions 253–4 double network hydrogels 149 double-quantum (2Q) dipolar recoupling 531 double reptation 24 double vacuum bag (DVB) process 135–6, 157 dough molding compounds (DMC) 929–31 DPD see dissipative particle dynamics drop test 847 droplet breakup 165–6, 168–74, 187–92, 201–5, 224–5, 318–22, 338 droplet coalescence 165–6, 174–7, 187–8, 214–18, 321–3, 338 droplet-matrix morphology 161–2, 254 DRS see dielectric relaxation spectroscopy drug delivery systems applications 895–6 interfacial properties 81, 90–2, 101–3 microgels 143–4 solid-state NMR spectroscopy 535, 543–4 DSC see differential scanning calorimetry DTP see dynamic plane source dual fillers 966–7

dual-laminated pipes 877 ductility 810, 815–16, 827 dynamic density functional theories (DDFT) 64–5 dynamic light scattering (DLS) 153, 639, 659–60 dynamic mechanical analysis (DMA) 9, 108–9 interfacial phenomena 500–1, 504–7 manufacturing techniques 133, 142 mixing and phase separation 487–9, 492–4 molecular dynamics 480, 500, 513 polymer blends 271–6 waste management 939 dynamic mechanical thermal analysis (DMTA) 271–6 dynamic phase behavior 647–51, 652–5, 729–30, 740–3 dynamic plane source (DTP) 394 dynamic polymer self-assembly 561–4 dynamic viscoelasticity 329–32, 340–5 dynamic vulcanization 274 dynamical models 62–5 E-glass 138, 874, 876–9, 907–8 EBCL see electron beam chemical lithography eccentric cylinder devices 221 eco-composites 297 ECRC see European Composites Recycling Company EDAX see energy dispersive analysis by X-ray Edwards Hamiltonian 35, 43 Edwards models 60 EDX see energy dispersive spectrometers effusivity, thermal 389, 393–4 EGA see evolved gas analysis Einstein equation 13, 327–8 Einstein parameters 431 elastic incoherent structure factor (EISF) 713 elastic moduli 85, 98 elastic recoil detection (ERD) 111 elastic scattering 594–5, 707–8, 710–13 elastically active chains (EAC) 812 elasticity see viscoelasticity elastomers 959–70 ageing processes 812, 838 approximately spherical particles 960–3 carbon nanotubes 966 controlled interfaces 967 dual fillers 966–7 glassy deformable ellipsoidal particles 963–4 layered fillers 964–5 magnetic particles 965 modeling and simulations 968–70 morphological characteristics 164, 166, 199, 218, 229 nanofillers 959–80 polyhedral oligomeric silsesquioxanes 965–6

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Index porous fillers 967 silicification and biosilicification 967 swelling processes 961–3 thermal analysis 377 waste management 942–3 electric moduli 484–5 electrical ageing 797, 835 electrical percolation threshold 414 electrical properties 425–77 ageing processes 797, 833–9 applications 457–8, 866, 887, 899, 907–12 carbon black 430, 432–5, 453–4, 457 carbon nanotubes 443–9, 507–12 characterization 9–10 conducting polymer-coated fillers 451 conducting polymers 9–10, 449–51, 530–1, 613–19 conductivity 425–72, 484–5, 496–8, 507–12, 866, 899 contact resistance 469–72 crystallinity 454–5 dielectric response 425–6 dielectric spectroscopy 484–5, 496–8, 507–12 fillers 426–7, 430–51, 455–6 four-probes method 460–5 graphite fillers 438–43 metallic fillers 435–7 metallized fillers 438 modeling and simulation 426–32 morphological characteristics 435, 456–7 multipercolation effect 456–7 percolation models 426, 427–32 polymer blends 428 polymer composite blends 449, 452–3, 456–7 polymer composites/nanocomposites 426, 432–72 polymer–filler interactions 455–6 processing conditions 452–7 resistance measurements 458–72 secondary processing 453 spreading resistance of contacts 468–70 two-probes method 459 Van der Pauw method 465–7 X-ray photoelectron spectroscopy 600, 613–19 electrical trees 835, 837 electromagnetic interference (EMI) 435–6, 438, 457 electromagnetic properties 965 electromechanical breakdown 834 electron affinities 450 electron beam chemical lithography (EBCL) 598 electron spin echo envelope modulation (ESEEM) 560, 562, 575, 579 electron spin echo (ESE) 560 electron spin resonance (ESR) spectroscopy 551–84

987

copolymers 560–9 crosslinking 571 dynamic polymer self-assembly 561–4 grafted copolymers 569 mechanical properties 287 morphological characteristics 8 nitroxyl radical examples 553–4 polymer blends 569–71 polymer composites/nanocomposites 557–8, 573–6 semi-interpenetrating polymer networks 570, 571–3 spin probe/spin label techniques 552, 557, 563–5, 574 techniques and principles 551–5, 558–60 theoretical background 555–60 electronic shearography (ES) 790–3 electronic speckle pattern interferometry (ESPI) 788–90 electronics industry applications 866, 907–12 electrostatic discharge (ESD) 457, 892–3 electrostatic interactions 85, 616–18 electrostatic separation 922–3 ellipsoid models 60–1 ellipsometry 109 elongation at break (Eb) 9, 260–4, 294, 936, 943 elongation processes 810–11, 813 elongational flow 218–21 elongational stress 453 Elshelby tensors 304–5 ELV see End-of-Life Vehicles embrittlement 811–12, 815–16, 823, 827 EMI see electromagnetic interference emission angle ERD 111 emulsions 330–5 encapsulated graphite (EG) 402 end-of-life criteria 798, 799, 921, 940 End-of-Life Vehicles (ELV) 921, 944 end-pinching 172, 224 end-splitting 202 endohedral filling of CNTs 449 energy dispersive analysis by X-ray (EDAX) 155 energy dispersive spectrometers (EDX) 587 energy industry applications 865, 877–8 energy recovery 921–3 energy transfer 706 engineering thermoplastics 276–7 entangled systems 14, 21–7 entanglement plateaux 340–1 environmental factors ageing processes 797, 798–9, 836–7, 839 applications 873–4 fire retardancy 843–4, 862 see also recycling; waste management

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environmental scanning electron microscopy (ESEM) 155 environmentally adaptive clothing 144 EOS see equation of state epiradiator test 847 epitaxial growth 90–1, 442–3 epoxidixed natural rubber (ENR) 127, 214 epoxy resins ageing processes 807, 823, 832, 838 applications 610–12, 867, 872–9, 882, 888, 890, 907–8 characterization 375, 609–12, 627, 650–1 chemical reactions 609–10 interfacial properties 98, 103 manufacturing techniques 134–6 mechanical properties 279–82, 300–3 synthesis 609 thermophysical properties 397, 404, 407 waste management 948–9, 951 equation of state (EOS) theory 42, 59–60, 647 ERD see elastic recoil detection ES see electronic shearography ESCA see X-ray photoelectron spectroscopy ESD see electrostatic discharge ESE see electron spin echo ESEEM see electron spin echo envelope modulation ESEM see environmental scanning electron microscopy ESPI see electronic speckle pattern interferometry ESR see electron spin resonance ethylene-butene copolymers 139 ethylene glycol dimethacrylate (EGDMA) 488 ethylene-1-hexene (EH) copolymers 124 ethylene-methacrylic acid (EMAA) copolymers 578 ethylene–octene copolymer (POE) 275 ethylene propylene diene monomer (EPDM) gas transport mechanisms 769 manufacturing techniques 126, 130 mechanical properties 263–4, 277 morphological characteristics 211, 225, 228 rheological properties 316 waste management 937–8 ethylene-propylene-maleic copolymer (EPM) 205 ethylene-propylene rubber (EPR) light scattering techniques 660–1 manufacturing techniques 128, 131–2 mechanical properties 277, 279 morphological characteristics 174, 183–4, 189, 192–5, 207, 214–17 ethylene-vinyl alcohol copolymer (EVOH) 225 ethylene vinyl acetate (EVA) applications 895, 902 electrical properties 457

fire retardancy 853, 855, 857–61 manufacturing techniques 130, 142 mechanical properties 253–4, 265, 269, 272–5 morphological characteristics 167, 186, 195, 208, 214, 229 rheological properties 329–30, 335 thermal analysis 379 thermophysical properties 403, 406–9 waste management 941 ethylene-vinyl alcohol copolymer (EVOH) 754, 756, 763–4 European Composites Industry Association (EuCIA) 921, 944 European Composites Recycling Company (ECRC) 921 European Directive 2000/53/EC 921 europium oxide 142 evaporation controlled kinetics 804–5 evolved gas analysis (EGA) 378–80 exfoliation characterization 9, 672–3 fillers 440, 442–3 future developments 11 mechanical properties 291 morphological characteristics 163–4, 178–82 polymer nanocomposites 2, 4 expandable interphases 103 expanded graphite (EG) 142, 440 extensional flow mixer (EFM) 220–1 extrusion future developments 11 interfacial properties 96, 99 manufacturing techniques 123, 128, 129–32 microfibrillar reinforced composites 671 morphological characteristics 162, 164–7, 181–2, 184–7, 207–8, 220–31 rheological properties 352–3 fast Fourier transforms (FFT) 656 fast scanning rate calorimetry (FSC) 365 feedstock recycling 928–9, 945–6 FEM see finite element model FENE see Finitely Extensible Nonlinear Elastic Fermi-pseudo potentials 709 ferroelectric properties 830 ferromagnetism, applications 911 Feynman-Kac formula 46 FFT see fast Fourier transforms FIB see focused ion beam fiber breakage 780, 785 fiber–matrix debonding 779 fiber-matrix morphology 161–2

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Index fiber metal laminates (FML) 886–7 fiber-reinforced composites (FRC) 10, 100 fiber-reinforced plastics (FRP) 921–2 fiberglass see glass fiber fibers ageing processes 831 applications 866–9, 872–4, 878–80 electrical properties 431, 451 interfacial properties 84 mechanical properties 253–4, 283, 294–7 solid-state NMR spectroscopy 546 thermophysical properties 405–6 waste management 924–5, 929–33, 935–40, 944–51 Fick’s laws 750–1, 804 field-based models 61–2, 64–5 field strength 833–5 filament winding process 139–40 filamentary breakdown 834–5 filler-matrix adhesion 302–4 fillers 2 ageing processes 826, 838 applications 865, 867–8, 907 carbon black 430, 432–5 carbon nanotubes 443–9, 966 compatibilization 455–6 conducting polymer-coated 451 electrical properties 426–7, 430–51, 452–3, 455–6 electrically conductive 432–51 fire retardancy 844, 849–50, 852–62 gas transport mechanisms 752, 757–63 graphenes 441–3 graphite 438–43 interfacial properties 95 life cycling 10 manufacturing/processing techniques 434–5, 439–42, 452–3 mechanical properties 253, 265–6, 298–304 metallic 435–7 metallized 438 rheological properties 14–27, 329–30 thermal analysis 361, 377 thermophysical properties 406–19 waste management 922, 923, 931–43, 951 see also nanofillers finite element model (FEM) 89, 415–19, 889 Finitely Extensible Nonlinear Elastic (FENE) potential 59 fire retardancy 843–64 combustion of polymers 844–5, 846–9 fillers 844, 849–50, 852–62 flame retardancy 845–6, 849–8, 877, 923, 926–7, 932

989

interfacial properties 95–6, 102 laboratory fire testing 846–9, 853–61 polymer composites/nanocomposites 844, 849–62 synergistic effects 854–61 fire triangle 844 flame retardancy 845–6, 849–62, 877, 923, 926–7, 932 flammability 846–8 flash method 390–1, 393 flash pyrolysis 379 flax-reinforced polypropylene 138 flexural moduli 300–1, 303 flexural properties 9 float and sink method 922 Flory exponent 36–7 Flory–Huggins parameter 40, 41–3, 86, 369, 371 Flory–Huggins theory 15, 31–2, 39–43, 67 analytical theories 51–3 binary blends 39–40 characterization 369, 371, 531–2, 647 de Gennes’ theory 41, 51–2 inhomogeneous systems 41 interfacial properties 85–7 real systems 41–3, 61 flow alignment 345–7 flow behavior see rheological properties flow curves 329–30, 349 flow field types 218–21 flow-induced morphological changes 345–7 flow instabilities 350–3 flow-light scattering 643–5 fluctuation effects 48–9 fluidized bed processing 929–31 fluorescein 144 fly ash 935 focused ion beam (FIB) analysis 111 forward recoil spectroscopy (FRES) 111 four-probes method 460–5 four-roll mills 219–20 Fourier transform infrared (FTIR) spectroscopy 8, 109–10, 142, 155, 378–80 Fourier transform neutron scattering 712–13 fractals 14 fractionation 927 fracture properties 260–1, 264, 284, 300–1 free radical polymerization 146 free volume theory 751 freely jointed chain model 36 freeze-drying 126 FRES see forward recoil spectroscopy Fresnel reflectivity 717, 736 friction coefficients 14–17

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FSC see fast scanning rate calorimetry FTIR see Fourier transform infrared fuel cells 458, 869, 909 full IPNs 150 functional graded materials (FGM) 900 functionalized CNTs 448 functionalized elastomers 164 functionalized graphenes 443 functionalized polymers 456, 600–2 fusion characteristics 9 g-methacryloxypropyl trimethoxy silane (g-MPS) 137 gas chromatography (GC) 378 gas transport mechanisms 749–75 active films 764–6 ageing processes 801, 802–3, 823 barrier multiphase materials 749, 752–67 definitions of transport parameters 749–51 dense polymer films 749–51 dispersions within polymer matrix 752–63, 769 filler size and orientation 755–60 impermeable spheres 752–7, 769 lamellar structures 752, 757–63, 769 multilayer systems 752, 763–5, 767 organic–inorganic membranes 767–9 polymer blends 755–7 polymer nanocomposites 752, 755–63 selective multiphase materials 749, 767–9 Gaussian chain models 34–5, 44, 47 GC see gas chromatography; glassy carbon generic modeling 57 geometrical percolation models 430 geometrical permeability models 757–63 Gibbs-Thomson equation 367 Ginzburg-Landau models 62, 65, 69–70 Ginzburg parameters 49 GISANS see grazing incidence small angle neutron scattering glass fiber-reinforced aluminum alloys (GLARE) 882–3 glass fiber-reinforced plastic (GFRP) 295–6, 873–4, 878–9, 882 glass fibers applications 871–4, 877–9, 881, 884, 890 manufacturing techniques 134, 137, 138, 139 waste management 926, 930–3, 937–9, 944–51 glass transition temperature 9 ageing processes 801, 805, 813, 829 dielectric spectroscopy 480–1, 490, 492, 497, 499–500, 506 electron spin resonance spectroscopy 558

interpenetrating polymer networks 284 modeling and simulations 38 polymer blends 313–14 solid-state NMR spectroscopy 528 thermal analysis 365–6, 371–2, 375–7 unfilled polymer systems 395 waste management 936 X-ray photoelectron spectroscopy 619–20, 624 glassy carbon (GC) plates 600 glassy polymers 970 global degradation rate 814 glycerol monomethacrylate (GMA) 208, 629 gold fillers 437 gold nanoparticles 104 gold-POEGMA overlayer thickness 592–7 graft copolymers characterization 530, 532, 543, 569, 589–90, 600–5 gas transport mechanisms 763 mechanical properties 261–3 waste management 934–5 graft IPNs 282 graphene oxide (GO) 442–3 graphenes 441–3 graphite fillers 438–43 graphite/epoxy composites 888–9 grazing incidence small angle neutron scattering (GISANS) 724–9, 733–43 Green Label scheme 921 Griffith’s criterion 834–5 ground fillers rubber from tires (GTR) 935 ground state dominance 53 guarded hot-plate method 390 guided tissue regeneration (GTR) 900 Gusev–Lusti model 759, 760 halogenated flame retardants 843, 850, 855, 927 halogenated olefin polymers 397 halogenation reactions 806–7 Halpin-Tsai equation 268–9, 288–9, 304–6 hand layup process 134 hardness 9 Hartmann-Hahn conditions 521, 526 Hashin–Shtrikman model 406 Hasselman–Johnson model 408 Hatta–Taya model 406 Havriliak-Negami (HN) expression 484, 490–1, 499, 503–4 heat capacity see specific heat capacity heat flux DSC 365 heat release rate (HRR) 846, 848–9, 855, 857–8 heat transfer 388

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Index heterophasic propylene-ethylene copolymers (HPEC) 565–6 hexabromocyclododecane (HBCD) 843, 850 HI see holographic interferometry hierarchic IPs 95 high-density polyethylene (HDPE) 132 characterization 526, 661, 672, 681–93 electrical properties 437, 451, 457 mechanical properties 261–5, 272–5 morphological characteristics 175–7, 184–6, 190–9, 208–14, 217, 227–31 thermophysical properties 416–18 waste management 936–7, 941 high impact polystyrene (HIPS) 573, 927 high permeability layers (HPL) high pressure mixing method (HPMM) 299 high resolution 13 C NMR 521, 523–4, 525–36 high resolution electron energy loss spectrometry (HREELS) 109 high structure aggregates 434 high-temperature shear deformation (HTSD) 277–8 high-temperature thermal protection (TPS) 891–2 hindered amine stabilizer (HAS) 565–7 HN see Havriliak-Negami Hoffman reactions 941 hole formation 225–6 hollow fibers 888 holographic interferometry (HI) 788–90 homopolymer blends 39–40, 51–2, 55–6, 65–7 honeycomb defects 786 honeycomb structures 136, 867, 882, 885, 891 hot radicals 845 hot-wire method 390 hot-wire parallel technique 394 HREELS see high resolution electron energy loss spectrometry humid ageing 797 humidity-induced recrystallization 90 humidity sensors 911–12 hyaluronic acid (HA) 900 hybrid carbon–glass fiber-reinforced plastic 873–4 hydrated fillers 849–50 hydrocarbon polymers 397–8 hydrodynamic modeling 13–14, 18 hydrogels applications 895–6, 900 interfacial properties 104–5 manufacturing techniques 148–9 mechanical properties 290–5 hydrogenolysis 928

hydrolytic ageing 826–9, 946 see also water-induced damage hydromagnesite 850 hydrotalcite 857–8, 860 hydroxyapatite (HA) 126, 897–8, 900, 904–5 hydroxycarbonates 850 2-hydroxyethyl methacrylate (HEMA) 537, 565 hydroxypropyl methacrylate (HPMA) 565 hydroxypropyl methyl cellulose (HPMC) 124–5 hygroscopic effects 889 hyperDSC 365 hyperfine interactions 555 ICP see inductively coupled plasma; inherently conducting polymers ideal chains 34–6 identification of waste materials 922–4 IEC see ion exchange chromatography IED see improvised explosive devices IGC see inverse gas chromatography immiscible polymer blends 2 impact properties 9 hollow fibers 888 nondestructive testing 779–80 polymer blends 276–9, 299 impact strength 938 impedance 389, 484–5 improvised explosive devices (IED) 880 in-flight monitoring 886 in-line electron beam treatment 96, 99, 112 in situ intercalative polymerization 142 in situ monitoring 11 incineration 921, 923, 929–32 incipient hole formation 225–6 inclusion 94–5 incoherent scattering function 711, 712–13 indium tin oxide (ITO) 892–3 inductively coupled plasma (ICP) 601 inelastic scattering 707–8 inert-PAMPS gels 149–50 infrared (IR) spectroscopy 142, 156, 922, 952 infrastructure applications 878–9 inherently conducting polymers (ICP) 9–10, 449–51, 530–1, 613–19, 909 inhomogeneous systems 41 injection molding 262, 453, 694–5 insulating interlayers 101 insulating polymers 9–10 insulin delivery/release systems 104–5 intensity–time correlation function (ITCF) 660 interacting chain models 36–8

991

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interactive hole formation 225–6 intercalation characterization 9 fillers 439–40 future developments 11 manufacturing techniques 141–3, 146 morphological characteristics 163–4, 178–82, 230 polymer nanocomposites 2, 4 X-ray scattering 672–3 interchain correlations 57 interchain scattering 731 interdiffusion 86, 94, 716, 719 interface analysis 105–11 interface distribution function (IDF) 678 interfacial agents 932–5 interfacial hydrolysis 833 interfacial properties 1–5, 81–121 basic characteristics 81–3 characterisics of interfacial layers 83–9 compatibilization 87, 93–100 controlled morphology 90–2 definitions and classification 82–3 dielectric spectroscopy 499–507 elastomers 967 electrical properties 430, 447 fire retardancy 859–60 future developments 111–12 homopolymer blends 66 interactions 84–5 interface analysis 105–11 kinetic factors 87–8 mechanical properties 262–3, 277 mobility 84, 92–3, 108–9 modeling and simulations 53 modifications to IPs 89–100 morphological characteristics 8, 175–6, 192–7 neutron scattering 733–43 phase border regions 106–7 reactivity 100–1 responsive interphases 101–5, 112 rheological properties 323, 335–7 structural properties 83, 88–9 structure–mechanics relationship 88–9 thermal analysis 377, 380 thermodynamic factors 85–7 interfacial tension 321 intermittent fibers 283 internal interfaces 66 interparticle distance (IPD) 277, 291 interpenetrating blends (IPB) 225 interpenetrating polymer networks (IPN) 3–5

future developments 11 interfacial properties 94 manufacturing techniques 149–57 mechanical properties 251, 252, 282–90 morphological characteristics 285–6 structure–property relationships 479–80, 482, 486–95 X-ray photoelectron spectroscopy 627–8 interphases (IP) characterization 83–4, 376–7, 576 compatibilization 87, 93–100 controlled morphology 90–2 definition 82–3 future developments 111–12 interface analysis 105–11 kinetic factors 87–8 mobility 84, 92–3, 108–9 modifications 89–100 non-reversibly adaptive 101–3 reactivity 100–1 responsive 101–5, 112 smart reversibly adaptive 103–5 structure–mechanics relationship 88–9 thermodynamic factors 85–7 intrinsic chemical stability 817–18 intumescent flame retarders (IFR) 857–8, 860 inverse gas chromatography (IGC) 108, 615 ion exchange chromatography (IEC) 617 ion exchange polymer-metal composites (IPMC) 890–1, 900–2 ion exchange X-ray photoelectron spectroscopy 616–18 ionic liquids 927 ionization potentials 450 ionomers 456 IPN see interpenetrating polymer network IR see infrared iron fillers 435–6 iron (III) oxide 853–4, 859 irreversible hydrolysis 826–8 isotactic polypropylene (iPP) 143, 360–4, 368–70, 377, 661 isothermal crystallization 361–3 isotropic scattering 653 isotropization 671, 690 Johnson-Kendall-Roberts (JKR) model 108 Kelly-Tyson model 937–8 keratin feather fibers (KF) 911 Kerner equation 269–70, 334–9 ketoprofen 544 Kiesseg oscillations 717

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Index Kim–Burns model 371–3 kinetic Monte Carlo (MC) methods 62 Kivelson theory 557 Kramers expression 18 Kramers-Kronig relation 426 Krieger-Dougherty equation 328–9, 331 Kuhn lengths 35, 41, 481 Kunori model 270, 288–9 L-TMA see localized thermomechanical analysis laboratory fire testing 846–9, 853–61 Lagrangian particle methods 224–5 Lambert-Beer expression 715 lamellar double hydroxides (LDH) 860–1 lamellar structures gas transport mechanisms 752, 757–63, 769 morphological characteristics 161–2, 166–8 neutron scattering 732 X-ray scattering 695 laminar mixing 452 laminar morphologies 253–4 laminated object manufacturing (LOM) 137 Langmuir-Blodgett (LB) films 83 large amplitude oscillatory shear flow (LAOS) 170–1, 345 laser flash analysis (LFA) 380 laser interferometry 788–90 latex blending 2, 125 latex IPNs (LIPN) 153–4 lattice-Boltzmann models 70 lattice chain models 36, 58–9 lattice cluster theory 42–3 layer-by-layer (LbL) assembly 911–12 layer thinning 225–6 layered double hydroxides 964–5 layered fillers 964–5 layered silicates see polymer-layered silicate nanocomposites LB see Langmuir-Blodgett LbL see layer-by-layer LCA see life-cycle assessment lead magnesium niobate-lead titanate (PMN-PT) 908 Leibler theory 51–2 Lennard-Jones potentials 59, 969 Lewis–Nielsen model 412, 415, 416 LFA see laser flash analysis life-cycle assessment (LCA) 10, 921, 940 Lifshitz points 49 light scattering techniques 8, 639–68 crystallization 660–72 experimental setup 642–3

993

flow-light scattering 643–5 historical development 639–40 instrumentation 644–5 intensity calibration 645–6 methodology and techniques 640–6 morphological characteristics 649–51, 656–8 multi-scale approaches 655–6 phase behavior of multiphase polymer systems 640, 646–56, 661 polymer blends 646–7, 651–8, 660–2 polymer gels 658–60 reaction-induced phase separation 649–51 rheology-light scattering 644–5 sample preparation 643 shear flow conditions 651–5 theoretical background 640–2 limiting oxygen index (LOI) 846–7, 853, 855, 861 linear correlation function (CF) 678, 680–1 linear elastic fracture mechanics 780–1 linear low-density polyethylene (LLDPE) light scattering techniques 660–1 mechanical properties 272 morphological characteristics 175, 198–9, 216, 220 rheological properties 325 liquid crystalline polymers (LCP) 648, 656, 925 liquid-to-glass transitions 800–2 literature review 5–7 living polymerization 112 localized thermomechanical analysis (L-TMA) 376–7 LOM see laminated object manufacturing London forces 613–15 long period 677 Lorentz corrected SAXS 677–8 Lorentz functions 713, 740–1 loss moduli 341–2, 346, 353 low shear rate steady state (LSRSS) viscosity 19–20 low-density polyethylene (LDPE) electrical properties 437, 451 fire retardancy 857–8 mechanical properties 275, 277–8 morphological characteristics 165, 209, 217, 225, 228–31 rheological properties 325 waste management 936 lower critical solution temperature (LCST) 104, 646–7, 895 lubricants 648 Lyapunov exponents 222–4 macromolecular interfacial properties 95–8, 100 macroscopic coarse-graining 33 macroscopic interfacial properties 82, 88–9

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macroscopic physical properties 8–10 macroscopic rheological properties 18, 24–5 magic angle spinning (MAS) 519, 521–2, 525–36 magnesium di-hydroxide (MDH) 849–50, 855, 859–61 magnetic properties 435–6, 438, 457, 965 maleic acid (MAH) 370 maleic anhydride (MA) dielectric spectroscopy 508–9 electrical properties 456 fire retardancy 860 gas transport mechanisms 763 manufacturing techniques 131 mechanical properties 265, 277, 279 morphological characteristics 178–9, 201–2, 205–8, 214, 228–9 waste management 934–5, 937, 939, 942–3 manufacturing techniques 123–60 batch/continuous mixers 128–9 cost effectiveness 123 electrical properties 434–5, 439–42, 452–7 extruders 128, 129–32 filament winding process 139–40 fillers 434–5, 439–42, 452–3 freeze-drying 126 future developments 156–7 hand layup process 134 in situ intercalative polymerization 142 interpenetrating polymer networks 149–57 latex blending 125 mechanical blending 126–32 mechano-chemical blending 132 melt intercalation/blending 127–8, 143 nanostructured gels 146 polymer blends 123–33 polymer composites/nanocomposites 123, 133–43, 157 polymer gels 143–9, 157 pultrusion 138–9 reaction injection molding 140, 157 resin transfer molding 136–8, 157 rotational molding 140–1 secondary processing 453 solution blending 124–5 solution intercalation 141–2 spray layup process 134–5 supercritical fluids 132–3 topological networks 146–7 vacuum bag molding 135–6, 157 Marangoni stress 176–7, 204 marine applications 871–5 MAS see magic angle spinning mass spectrometry (MS) 378 mass transfer induced hydraulic actuators (MTIHA) 901

mass transit applications 880–2 matching lattice size theory 368 material ageing see ageing processes matrix cracking 779–83, 786 matrix deformations 779 matrix-dispersed structures 208–9 MAXS see middle-angle X-ray scattering Maxwell equation/law 405–8, 416, 753–6, 768–9 Maxwell model for viscoelasticity 18, 27 Maxwell relaxation time ratio 189–90, 216 Maxwell-Wagner-Sillars (MWS) polarization/relaxation 502, 509 MC see Monte Carlo MD see molecular dynamics MDSC see modulated DSC mean field approaches 47 mean geometric model 412–13 mechanical ageing 798, 799 mechanical blending 126–32 mechanical compatibilization 100, 102 mechanical properties 251–310 alloys 251 characterization 9 compatibilization 260–5 compositional effects 257–60, 273 crosslinking 260 dynamic 271–6 fillers 253, 265–6, 298–304 future developments 307 impact properties 276–9, 299 interfacial properties 85, 88–9, 100 interpenetrating polymer networks 251, 252, 282–90 manufacturing techniques 140 mixing conditions 256–7 modeling and simulations 266–71, 275–6, 287–90, 304–6 morphological characteristics 253–6, 270–4, 285–6 polymer blend nanocomposites 252, 279–81 polymer blends 251–2, 253–81 polymer composites/nanocomposites 251, 253, 293–306 polymer gels 251, 252, 290–5 viscoelasticity 251, 269, 272 waste management 929–31, 934–43, 949–51 see also rheological properties mechanical recycling 932–43, 947–51 mechanical relaxation 343–5 mechanical separation of waste 924 mechano-chemical blending 132 medical applications 866, 895–907 see also drug delivery systems

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Index melamine phosphates 851–2, 855, 857 melt blending 671 melt intercalation/blending 127–8, 143 melt mixing 2, 178–82, 215, 257–8 melting temperatures 8–9 membranes 767–9, 909 MEMS see microelectromechanical systems mercury porosimetry 903–4 Meredith–Tobias model 406–7 mesoscopic coarse-graining 33–4 metal recycling 924 metallic fillers 435–7 metallic matrix composites (MMC) 778 metallized fillers 438 methacrylic polymer (PM) 571–2 2-methacryloyloxy phosphorylcholine (MPC) 629 method of images 463 methyl acrylate (MA) 275–6 micellar structures 726–8, 730, 733–43 microelectromechanical systems (MEMS) 885–6, 901 microelectronic sensors 144 microfibrillar reinforced composites (MFC) 670, 671–2, 678, 681–93 microgels 143–5 micromechanical cleavage 442 micromechanical testing 108 microphase separation 312, 339–40, 343, 565 microporosity 434 micro-Raman 109–10, 143 microscopic coarse-graining 33 microscopic interfacial properties 109 microscopic order with macroscopic disorder (MOMD) 558 microscopic rheological properties 15–17, 22–4 microsensors 885–6 microspheres 605–8 microthermal analysis 376–7 middle-angle X-ray scattering (MAXS) 674 miscible polymer blends 2 Mitsui Silicon Containing Polymer (MSP) 891–2 mixing conditions 256–7 mixing sequence effects 227–9 mobility 84, 92–3, 108–9 modeling and simulations 7–8, 31–80 ABC miktoarm star copolymers/selective solvent 38 analytical theories 49–55 applications 55–6, 65–70 atomistic models 58 basic concepts of polymer theory 32–9 binary blends 39–40, 51–2, 55–6, 65–7

chain dynamics 38–9 coarse-grained models 33–4, 56–61, 70 concentrated polymer solutions 37 copolymer conformations 53–5 de Gennes’ theory 15–16, 41, 51–2 dense melts 38 dynamical models 62–5 elastomers 968–70 electrical properties 426–32 field-based models 61–2, 64–5 first order models 404–5 Flory-Huggins theory 15, 31–2, 39–43, 51–3, 61 fluctuation effects 48–9 fundamental properties of polymer molecules 32 future developments 70 gas transport mechanisms 757–65, 769–70 Gaussian chain model 34–5, 44, 47 generic modeling 57 ground state dominance 53 ideal chains 34–6 inhomogeneous systems 41 interacting chain models 36–8 interfacial profiles 53 Leibler theory 51–2 mean field approaches 47 mechanical properties 266–71, 275–6, 287–90, 304–6 multiphase polymer systems 39–70 multiscale modeling 8, 32 nanofillers 968–70 numerical prediction models 415–19 Ohta-Kawasaki Functional 51–2 particle-based dynamics 62–4 polymers in solution and blobs 36–8 random phase approximation 32, 50–2 rheological properties 13–29, 326–39 scaling behavior 36–7 second order models 405–10 self-consistent field theory 7, 31–2, 43–9, 68 semi-empirical prediction models 412–15 strong segregation theory 32, 41, 46, 52–5 strong stretching theory 53–5 structural models 58–62 systematic bottom-up modeling 57–8 theoretical aspects 31–56 thermophysical properties 395, 404–19 third and fourth order models 410–12 weak segregation limit 41, 46, 50–2, 55–6 modulated DSC (MDSC) 9 molar absorptivity 807–8 molar mass 810–12, 828–9 molecular density functionals 62

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molecular dynamics (MD) 62 dielectric spectroscopy 480, 482, 499–500, 513 electron spin resonance spectroscopy 574 neutron scattering 729–30, 740–3 molecular sieves 767–9 MOMD see microscopic order with macroscopic disorder momentum transfer 706 monochromators 718–19, 721–3 monomer incompatibility 42 monomethoxylated polyethylene glycol (mmePEG) 564 Monte Carlo (MC) methods 58, 62, 69, 968–9 montmorillonite (MMT) 103–4 elastomers 965 fire retardancy 853, 861 mechanical properties 267, 298–9, 302, 305–6 thermal analysis 367–8 X-ray photoelectron spectroscopy 626–7 X-ray scattering 678, 679–81, 695–7 see also organomontmorillonite Mori–Tanaka mean field theory 304–6 morphological characteristics 8, 161–249 ageing processes 797, 800–2, 829–31 block copolymers 339–40 chaotic mixing 221–7 compatibilization 197–208, 213–14 compositional effects 208–12 development mechanisms during processing 164–82 dielectric spectroscopy 480, 482, 499 droplet breakup and deformation 165–6, 168–74, 187–92, 201–5, 224–5 droplet coalescence 165–6, 174–7, 187–8, 214–18 electrical properties 435, 456–7 flow field types 218–21 gas transport mechanisms 751, 752, 757–63, 769 initial development mechanisms 164–8 intercalation, exfoliation and dispersion 163–4, 178–82, 230 interfacial properties 175–6, 192–7 interpenetrating polymer networks 285–6 light scattering techniques 649–51, 656–8 material-relevant factors 183–218 mechanical properties 253–6, 270–4, 285–6 mechanisms of blending 162 mixing sequence effects 227–9 polymer blend nanocomposites 162–4, 178–82, 187, 207–8, 212–18, 225–31 polymer blends 161–231, 253–6, 270–4, 311–18, 323–5, 334–9 processing parameters 230–1 processing-relevant factors 218–31 rheological properties 311–18, 323–5, 334–40

solid-state NMR spectroscopy 537 ternary and quaternary blends 205–7, 209–12 viscoelasticity 171, 174–6, 187–92, 219 viscosity 171, 175, 183–7, 214, 219 waste management 943 X-ray scattering 694–5 MS see mass spectrometry MTGA see temperature-modulated TGA MTIHA see mass transfer induced hydraulic actuators multi-frequency TMDSC 374–5 multilayer gas transport mechanisms 752, 763–5, 767 multimandrel filament winding 877 multipercolation effect 456–7 multiple crazing 277 multiscale modeling 8, 32, 655–6 multiwall carbon nanotubes (MWNT) applications 908, 910 characterization 508–12, 561–2 elastomers 966 electrical properties 444, 446–9 fire retardancy 853, 859 mechanical properties 300–1 morphological characteristics 218, 220, 229, 231 multiwall systems 894 musculoskeletal applications 895–905 MWS see Maxwell-Wagner-Sillars N-succinimidyl ester-functionalized pyrrole (PyNSE) 618–19 n-type conductivity 450–1 N-isopropylacrylamide (NIPAM) 145–6, 148–9, 151, 658–60, 895–6 N-methacryloyl-L-histidine (MHist) 148 nanoclays see clay nanocomposites nanocompatibilization 87, 94 nanocomposites see polymer nanocomposites nanofillers 959–80 approximately spherical particles 960–3 carbon nanotubes 966 controlled interfaces 967 crystallinity 970–1 dual fillers 966–7 elastomers 959–70 fire retardancy 844, 852–4, 856–62 glassy deformable ellipsoidal particles 963–4 glassy polymers 970 layered 964–5 life cycling 10 magnetic particles 965 modeling and simulations 968–70 naturally-occurring polymers 971–2

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Index polyhedral oligomeric silsesquioxanes 965–6 porous 967 rigid polymers 972 silicification and biosilicification 967 thermosets 972–3 nanoplatelets 11 nanosized carbon particles (NCP) 507–9 nanostructured blends see polymer blend nanocomposites nanostructured gels 146 nanothermal analysis 376–7 natural graphite (NG) 142 natural rubber (NR) characterization 530, 557–9, 571, 575 elastomers 964–5 electrical properties 456–7 manufacturing techniques 125, 126–7, 151–2 mechanical properties 257–8, 265–6, 270 morphological characteristics 214 waste management 932 naturally-occurring polymers 971–2 NDT see nondestructive testing near surface crystallization of micelles 733–43 necking 172, 224 negative-deviation blends (NDB) 185–6 neutron detectors 719 neutron reflectometry (NR) 713, 714–16, 719–20, 724–5, 733–43 neutron scattering 8, 705–47 Born approximation 708–10, 713, 716, 734–5 contrast variation 730 copolymers 730, 731–3 crystallinity 726–9, 733–43 elastic/quasielastic scattering 707–8, 710–13 experimental results 729–43 instrumentation 719–24 micellar structures 726–8, 730, 733–43 near surface crystallization of micelles 733–43 polymer blends 729–31 production and detection of neutrons 716–19 properties of neutrons 705–6 shear thinning/dynamics 730–3, 739–43 small momentum transfer 713–16 techniques and principles 706–8, 713–29, 730 theoretical background 705–16 time and length scales 707 neutron spin echo (NSE) method 723–4 Newman projections 524–5 nickel fillers 436–7 Nielsen’s law 430–2, 756, 757–60, 763 Nishi–Wang plots 369–70, 373 nitrile rubber (NBR) 125, 253–5, 269

997

nitrogen tetroxide (NTO) 893 nitroxide mediated polymerization (NMP) 600 NMR see nuclear magnetic resonance N,N’-dicyclohexyl-2,6-naphthalene-dicarboxamide (NJS) 361, 362 noncrimp fabric (NCF) 866 nondestructive testing (NDT) 10, 777–96 acoustic emission 782–3 electronic shearography 790–3 failure mechanisms 779–81 laser interferometry 788–90 optical deformation and strain measurement 791–5 polymer composites 777–96 radiography 785–6 techniques and principles 777–9 thermography 786–8 ultrasonic scanning 783–5 visual inspection 781–2 non-reversibly adaptive interphases 101–3 non-swollen gels 5 nonwoven fabric 879–80 NR see natural rubber; neutron reflectometry NRA see nuclear reaction analysis NSE see neutron spin echo nuclear magnetic resonance (NMR) spectroscopy manufacturing techniques 156 molecular dynamics 480 morphological characteristics 8 relaxation studies 538–44 spin diffusion 544–6 theoretical background 520–2 see also solid-state NMR spectroscopy nuclear reaction analysis (NRA) 111 nucleating agents (NA) 360–2, 368 nucleation and growth 648–9 numerical prediction models 415–19 nylon blends manufacturing techniques 131, 133 mechanical properties 254–5, 265–8, 279, 306 waste management 938 object grating method 791–5 OCLV see optimized compaction, low void off-lattice chain models 59–60 offshore applications 874–5 ohmic contacts 469–70 Ohta-Kawasaki Functional 51–2 Oldroyd model 331–4 OM see optical microscopy on-line visualization 165, 170, 174, 190, 317 OOT see order–order transition optical activity applications 895–6

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optical deformation and strain 791–5 optical fiber sensors 893 optical light microscopy 8 optical microscopy (OM) 650, 656 optical phase extraction 790 optical shear flow cells 171 optical transition energies 450 optimized compaction, low void (OCLV) 875–6 order–disorder transitions (ODT) 48–9, 68, 339 order–order transition (OOT) 345–6 ordered microphases 253–4 organic–inorganic hybrid materials 902 organic–inorganic membranes 767–9 organomodified lamellar silicates (OMLS) 856, 859, 862 organomontmorillonite (OMMT) fire retardancy 857–9, 861 manufacturing techniques 143 mechanical properties 275, 298–9, 302 waste management 934, 940–1 X-ray photoelectron spectroscopy 626–7 oriented polymer blends 755–7 orthotropic composites 390 osmotic cracking 832–3 Ostwald-Buz´agh continuity principle 93 Oswald ripening 736 oxazolidine derivatives 554 oxidation induction time (OIT) 825 oxidation reactions 806–8, 812, 814–15, 818–24, 839, 934 oxygen permeability 752 oxygen transport 823 p-type conductivity 450–1 packaging applications 764 packing volume fraction 408–12, 415, 431 paclitaxel 535, 544 Pal models 408–10 PALF-reinforced polyester composites 296–7 Palierne model 333–5 palladium fillers 437 parallel breakup 172 parallel imaging 598 parallel model 267–71, 275–6, 288–9, 404–5, 413–14 parallel plate devices 168–9 partial discharges 835 partial encapsulation 193 particle-based dynamics 62–4 particle size distribution (PSD) 277 partly swollen gels 5 PCB see printed circuit boards PCT see percolation-to-cluster transitions

PEFC see polymer electrolyte fuel cell PELDOR see pulsed double electron–electron resonance penetration depth 715 pentaerythritol 851–2, 857 percolation models 426, 427–32 percolation phenomena 507–12 percolation-to-cluster transitions (PCT) 649 permeability 9, 749, 751, 753–7, 760–7 permittivity 92–3, 426 peroxyl radicals 819–25 phase angle plots 342 phase border regions 106–7 phase diagrams 40, 47–9, 640, 652 phase orientation studies 676–7 phase separation 647–51, 655–6, 661 phase-transition properties 9, 10 phenolic resins 398 phenyl-tert-butyl-nitrone (PBN) 552 phosphorated melamine 851–2, 855, 857 phosphorus-based flame retardants 850–2, 855, 857–8, 861–2 phosphorylated epoxy resin (PEP) 95 photoacoustic techniques 393 photochemical ageing 797, 812, 815, 818–19, 821–3, 839 photogrammetry 791–5 photoionization 586–7 photolysis 100 photopolymerization 906 photopyrotechnic techniques 393 photostability 102 physical ageing 798, 799–806 additive migration 803–6, 826, 829 application-related factors 832–9 electrical insulations 833–9 multiphase polymer systems 837–9 solvent absorption 802–3 structural reorganization 800–2 water-induced damage 826–9, 832–3, 836–7 physical characteristics 1–5, 8–10 physical compatibilization 93–100, 102 PI see polydispersity index; polyisoprene piezoelectric properties 830, 897 piperidine derivatives 553 plastic deformations 779 plastic laminates 866 plasticization 577, 751, 802–6 plastics production and consumption 5–6 platelet formation 163–4, 180–2 plerospheres 894 PLM see polarized light microscopy

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Index PLS see positron lifetime spectrometry Pluronics 561–2, 733, 738, 741, 900 POEGMA overlayer thickness 592–7 polarized light microscopy (PLM) 130, 146, 359–60, 693–5 poly(α methyl styrene) (PαMS) 817 poly-trans-acetylene 450 poly(acrylamide) (PAM) 155, 658 poly(2-acrylamido-2-methylpropanesulfonic acid) (PAMPS) 149–50 poly(acrylic acid) (PAAc) 151 polyacrylonitrile (PAN) 150, 275–6, 293, 362, 565 poly(alkyl acrylates) (PAA) 486–99, 544, 578 polyamides (PA) ageing processes 810–11, 823, 828 dielectric spectroscopy 510–11 electrical properties 438 gas transport mechanisms 756, 761, 763–4 manufacturing techniques 128–9, 132 mechanical properties 277, 305 microfibrillar reinforced composites 671–2, 678, 679–93 morphological characteristics 162, 168, 175–6, 181–4, 187, 189, 193, 199–205, 211, 214–18, 225–31 rheological properties 322–3 solid-state NMR spectroscopy 543 thermal analysis 362, 369, 377, 379 thermophysical properties 396, 398, 402 waste management 922–3, 937–8, 943 poly(anhydrides) 898 polyaniline (PANI) 450, 613, 909–10, 912, 965 poly(anilinesulfonic acid) (SPANI) 911–12 poly[bis(trifluoroethoxy)phosphazene] (PBFP) 537–8 polybromo diphenyl ethers (PBDE) 843–4, 850 polybutadiene (PB) copolymers 564–5, 567, 653–4, 730, 819, 823–5 poly(butyl acrylate) (PBA) 457, 486–95 poly(butyl methacrylate) (PBMA) 486–93, 569–70 poly(butylene adipate) (PBA) 527–30 poly(butylene terephthalate) (PBT) 128–9, 214–17, 229, 265, 527–30, 826 poly(caprolactone) (PCL) applications 903 electron spin resonance spectroscopy 564, 573, 575 interfacial properties 100 light scattering techniques 661–2 manufacturing techniques 133, 143 mechanical properties 257–60, 277–8 rheological properties 313–14 solid-state NMR spectroscopy 526–7 X-ray photoelectron spectroscopy 588–90

999

polycarbonate (PC) ageing processes 801, 807, 823, 826 interfacial properties 95 light scattering techniques 661 manufacturing techniques 127–8, 131–2 morphological characteristics 165, 184, 205–6, 214 rheological properties 316 thermophysical properties 396, 404 waste management 936, 938, 941 poly(cis-butadiene) rubber (PcBR) 657 poly(cyclohexyl acrylate) (PCHA) 564, 570–1 poly(cyclohexyl methacrylate) (PCHMA) 570–1 poly(dimethyl acrylate) (PDMA) 621–3 poly(2,6-dimethyl oxyphenylene) 802 poly(dimethyl siloxane) (PDMS) ageing processes 823 applications 908 characterization 500–7, 629 elastomers 960–6, 969–70 manufacturing techniques 156 morphological characteristics 171, 174, 184, 190, 202–3, 219, 221, 229 poly(2-(dimethylamino) ethylmethacrylate) (PDMAEMA) 604–5 poly(dioxolane) 542 polydispersity index (PI) 809 polyesters applications 872–3, 881 manufacturing techniques 134–5, 139 solid-state NMR spectroscopy 526–30 thermophysical properties 396 waste management 949 poly(4-ethenylphenolmethylsiloxane) (PEPS) 621–3 poly(ether-block-amide) copolymers 542 poly(ether-ester) blends 256–7 poly(ether ether ketone) (PEEK) 823, 898–9, 904–7 poly(ether imide) (PEI) 661 polyethers 396 poly(ether sulfone) (PES) 650–1, 823 poly(ether-urethane) ionomer (PEUI) 154 poly(ethyl acrylate) (PEA) 492, 530 poly(ethyl methacrylate) (PEMA) 492, 496–8 polyethylene (PE) ageing processes 807, 810–11, 813, 817–19, 823–5 characterization 546, 588 electrical properties 456–7 gas transport mechanisms 762–3 manufacturing techniques 139 mechanical properties 271 morphological characteristics 174–7, 185–6, 192–3, 198, 208, 228

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polyethylene (PE) (Continued) rheological properties 317–18 waste management 941 poly(ethylene-co-butene) (PEB) 661, 673, 697–9 poly(ethylene-co-hexene) (PEH) 661 poly(ethylene dioxy)thiophene (PEDOT) 613 poly(ethylene glycol) diacrylate (PEGDA) 153, 531 poly(ethylene glycol) (PEG) characterization 540–1, 578, 590, 661–2 manufacturing techniques 124, 146–7, 156 morphological characteristics 184 poly(ethylene naphthalate) 830 poly(ethylene oxide) (PEO) electron spin resonance spectroscopy 561–3, 569, 576, 578 light scattering techniques 661 manufacturing techniques 124–5, 149 mechanical properties 282 neutron scattering 729–30, 732–3 rheological properties 26, 315, 335–8 solid-state NMR spectroscopy 532, 534, 544–5 poly(ethylene succinate) (PES) 124 poly(ethylene terephthalate) (PET) ageing processes 807, 817, 826, 830 gas transport mechanisms 754, 756 light scattering techniques 656, 660–1 morphological characteristics 183–4, 199–200 rheological properties 322–3 solid-state NMR spectroscopy 530, 542 thermal analysis 362, 365–6 waste management 928–9, 933, 936, 941, 943 X-ray scattering 673, 697–9 polyglycidol (PGL) 605–8 poly(glycidyl methacrylate) (PGMA) 601–2, 625–6 polyhedral oligomeric silsesquioxanes (POSS) 301–2, 853, 965–6 poly(hexylene oxide) (PHO) 282 polyhydroxyalkanoates (PHA) 530 poly-4-hydroxybutyrate (P4HB) 897 poly(3-hydroxybutyrate) (PHB) 124, 530 poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) 530 poly(hydroxyethyl acrylate) (PHEA) 487, 496–8 poly(2-hydroxyethyl methacrylate) (PHEMA) 545, 565, 597, 603–4 polyimides (PI) 767, 823 polyisobutylene (PIB) 190, 202–3, 219, 221, 545, 817 polyisoprene (PIP) 47, 343, 819, 823 polyketones 396 poly(L-lactic acid) (PLLA) 143, 526–7, 903, 939–40 poly(L-lysine) (PLL) 533, 536, 540

poly(lactic acid) (PLA) 257–60, 277–8 poly(lactic-co-glycolic acid) (PLGA) 126, 589–90, 900 poly(lactide) (PLA) 149 polymaleic anhydride octyl vinyl ether (PMAOVE) 577 polymer blend nanocomposites mechanical properties 252, 265–7, 279–81 morphological characteristics 162–4, 178–82, 187, 207–8, 212–18, 225–31 polymer blends 2 analytical theories 49–56 applications 65–70 blending laws and viscoelasticity models 326–35 coarse-grained models 56–8 development mechanisms during processing 164–82 dispersed phase morphology 3 droplet breakup and deformation 165–6, 168–74, 187–92, 201–5, 224–5 droplet coalescence mechanisms 165–6, 174–7, 187–8, 214–18 electrical properties 428, 449, 452–3, 456–7 electron spin resonance spectroscopy 569–71 Flory-Huggins theory 39–43 future developments 10–11 gas transport mechanisms 755–7 initial development mechanisms 164–8 intercalation, exfoliation and dispersion 163–4, 178–82, 200, 230 interfacial properties 86, 94 light scattering techniques 646–7, 651–8, 660–2 literature review 6–7 low frequency viscoelastic behavior 335–9 manufacturing techniques 123–33 material-relevant factors 183–218 mechanical properties 251–2, 253–81 morphological characteristics 161–231, 253–6, 270–4, 311–18, 323–5, 334–9 neutron scattering 729–31 oriented 755–7 processing-relevant factors 218–31 rheological properties 311–18, 322–39 self-consistent field theory 44–9 solid-state NMR spectroscopy 531–5, 539–42 specificity of blend rheology 322–6 structure–property relationships 479–80 swelling 325–6 thermal analysis 370–3 thermophysical properties 396–400, 419 viscoelasticity 316, 317, 326–39 waste management 932, 936–7 X-ray photoelectron spectroscopy 619–24 X-ray scattering 670, 693–5

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Index polymer composites 2 acoustic emission 782–3 ageing processes 838–9 applications 865–920 compatibilization 455–6 crystallinity 454–5 electrical properties 426, 432–72 electron spin resonance spectroscopy 557–8, 573–5 electronic shearography 790–3 failure mechanisms 779–81 fire retardancy 844, 849–52, 854–62 interfacial properties 81, 94 laser interferometry 788–90 literature review 5–7 manufacturing techniques 123, 133–41, 157 mechanical properties 251, 253, 293–7 multipercolation effect 456–7 nondestructive testing 777–96 optical deformation and strain measurement 791–5 production and consumption 5–6 radiography 785–6 resistance measurements 458–72 solid-state NMR spectroscopy 543 structure–property relationships 479–80, 482, 499–503 thermal analysis 375 thermography 786–8 thermophysical properties 387–8, 390, 394–420 ultrasonic scanning 783–5 visual inspection 781–2 waste management 921–57 X-ray photoelectron spectroscopy 624–7 X-ray scattering 670, 671–2, 681–93 polymer conformation 522–5 polymer electrolyte fuel cell (PEFC) 909 polymer gels 3 applications 895–6, 900 future developments 12 inhomogeneities 658–9 light scattering techniques 658–60 literature review 6–7 manufacturing techniques 143–9, 157 mechanical properties 251, 252, 290–5 sol–gel transition 659–60 swelling behavior 3, 5, 12 polymer melts 27 polymer microcomposites 298 polymer nanocomposites compatibilization 455–6 crystallinity 454–5 dielectric spectroscopy 479–80, 482, 499–507

1001

electrical properties 426, 432–72 electron spin resonance spectroscopy 575–6 fire retardancy 844, 852–4, 856–62 future developments 11 gas transport mechanisms 752, 755–63 interfacial properties 81, 94 literature review 5–7 manufacturing techniques 141–3, 157 mechanical properties 298–306 morphologies 2–4 multipercolation effect 456–7 preparation 2, 4 production and consumption 5–6 resistance measurements 458–72 rheological properties 14 solid-state NMR spectroscopy 538, 543 waste management 940–3 X-ray photoelectron spectroscopy 624–7 X-ray scattering 670, 672–3, 678–81, 683, 695–9 polymer networks see elastomers; interpenetrating polymer networks; thermosets polymer–clay nanocomposites (PCN) fire retardancy 857 mechanical properties 265–6, 305–6 morphological characteristics 163–4, 178–82, 207–8, 212–18, 220, 226–31 X-ray scattering 672–3, 678–81, 695–9 polymer–layered nanocomposites 141–2 polymer–layered silicate nanocomposites (PLSN) characterization 499–507, 538 mechanical properties 303–4 morphological characteristics 163–4, 178–82, 207–8, 212–18, 226–31 waste management 940–1, 943 polymeric matrix composites (PMC) 778 poly(methacrylic acid) (PMAA) 532, 569, 573, 600, 602 poly(methyl acrylate) (PMA) characterization 486–99, 564, 567, 575, 577 electrical properties 457 manufacturing techniques 142 poly(methyl methacrylate) (PMMA) ageing processes 803, 817, 823 applications 904 dielectric spectroscopy 493–5 electrical properties 457 electron spin resonance spectroscopy 564, 569–70, 576–7 fire retardancy 853–4, 859 interfacial properties 95 manufacturing techniques 125, 132, 156 mechanical properties 257–8, 280–1, 284–90, 293

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poly(methyl methacrylate) (PMMA) (Continued) morphological characteristics 162, 176, 188, 192, 195–8, 201–2, 205–6, 211–14, 219–20, 226 neutron scattering 729–30 rheological properties 323, 333 solid-state NMR spectroscopy 542 thermal analysis 372–3 waste management 932 X-ray photoelectron spectroscopy 595, 601–4, 615–16, 620–2 poly(methyl vinyl oxazoline) (PSox) 179, 202 poly[methylene (polyphenyl isocyanate)] (PMPPIC) 935 poly(methylphenylsiloxane) 649 poly(N-isopropyl acryl amide) 542 poly(N-isopropylacrylamide) 104 poly(N-isopropylacrylamide) (PNIPA) characterization 569, 577–8, 658–60 manufacturing techniques 146, 148–9, 151, 157 mechanical properties 290–2 poly(N,N-dimethylacrylamide) (PDMAA) 290, 292, 659–60 polyolefins 526, 763–4 poly(oxy methylene) (POM) 130–1, 183, 323, 456–7, 817–18, 823 poly(p-dioxanone) (PPDO) 124 poly(p-phenylene) 450 polypeptides 536 poly(phenylene sulphide) (PPS) 831, 938 poly(propionylethylenimine-co-ethylenimine) (PPEI-EI) 448 poly(propylene carbonate) (PPC) 569–70 poly(propylene fumarate) 898 poly(propylene oxide ) (PPO) 534, 545, 561–3, 730, 732–3 polypropylene (PP) ageing processes 807, 819, 823–4 applications 875, 903–4 dielectric spectroscopy 508–9, 512 electrical properties 455, 457 fire retardancy 858, 860 interfacial properties 90, 95 light scattering techniques 656–7, 660–1 manufacturing techniques 124, 127–32, 138–40 mechanical properties 254–6, 261–5, 272–5, 279, 296, 303–4 morphological characteristics 165, 168, 172–4, 178–9, 184–6, 189, 192–200, 207–8, 214, 217–18, 225–8 rheological properties 316, 335 thermal analysis 360–4, 367–70, 377 thermophysical properties 402, 404, 412–13

waste management 922–3, 925–6, 931–6, 939–40, 949–50 X-ray photoelectron spectroscopy 588, 613–15 X-ray scattering 670, 693–7 polypyrrole (PPy) 450–1, 613–19, 621, 626 polyrotaxanes 146–7 polysaccharides 398–9, 523 polysiloxanes 96, 98, 102, 104, 399 polysilsesquioxanes 146 polystyrene (PS) ageing processes 807 elastomers 963–4, 970 electrical properties 437 electron spin resonance spectroscopy 564–5, 567, 575 fire retardancy 857 interfacial properties 95 light scattering techniques 649, 652–4, 656 manufacturing techniques 128, 132, 151–4 mechanical properties 254–7, 271, 277–8, 280 modeling and simulations 42, 47 morphological characteristics 162, 165, 167, 172–9, 188–201, 205–6, 209–14, 219, 225 neutron scattering 730–1 rheological properties 313–14, 317–18, 323, 333, 343 thermal analysis 372–3, 379 thermophysical properties 414 waste management 927, 937 X-ray photoelectron spectroscopy 588–9, 605–8, 620 X-ray scattering 670, 693–5 polystyrene-block-polybutadiene-block-poly-(methyl methacrylate) (SBM) 281–2 poly(styrene-co-methacrylic acid) (STMAA) 570 poly(styrene-co-methyl methacrylate) (SMMA) 647–8 poly(styrene-co-4-vinyl phenol) (STVPh) 569–70 polysulfides 399 polysulfones (PSU) 132–3, 399 poly(tetrafluoroethylene) (PTFE) 416, 417–19, 569, 817, 823, 909–10 poly(tetrahydrofuran) 542 polythiophenes 450, 613 polyurea 571 polyurethanes (PU) applications 898 electron spin resonance spectroscopy 567, 571–3 manufacturing techniques 142, 153, 155 mechanical properties 284–5 thermophysical properties 399 waste management 928, 935 poly(vinyl acetate) (PVAc) 125 poly(vinyl acetate-co-vinyl alcohol) (PVA-VA) 448

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Index poly(vinyl alcohol) (PVA) 155, 763, 766 poly(vinyl benzene) (PVB) 605–8 poly(vinyl chloride) (PVC) ageing processes 803–6, 807, 817–18 applications 871–2 characterization 372–3, 620–1 electrical properties 457 morphological characteristics 175, 198–9 rheological properties 314, 343 waste management 924–5, 934, 936, 939 poly(vinyl ethylether) (PVEE) 620 poly(vinyl isobutyl ether) (PVIBE) 533, 535, 540 poly(vinyl methyl ether) (PVME) 42, 652–4 poly(vinylidene fluoride) (PVDF) ageing processes 830 mechanical properties 265–8, 293 morphological characteristics 162 rheological properties 315, 335–8 thermal analysis 362, 364, 368–70 poly(vinylphenol) (PVPh) 532–3, 534, 542 poly(4-vinylpyridine) (PVPy) 621–4 poly(vinylpyrrolidione) (PVPr) 621–3 poly(vinylpyrrolidone) (PVP) 88, 532–3, 534, 542 Porod’s law 642 porosity 434 porous fillers 967 position-sensitive detectors (PSD) 724 positive-deviation blends (PDB) 185–6 positive-negative-deviation blends (PNDB) 185–6 positron lifetime spectrometry (PLS) 110–11 power compensated DSC 365 press forming 868–70 pressure vessels 876 printed circuit boards (PCB) 839, 907–8 probability distributions 22–3, 35, 224 processing agents 803 processing techniques see manufacturing techniques prosthetic applications 895–905 proteins 603–4 PSD see particle size distribution; position-sensitive detectors pseudo-IPNs 156 PtBMA grafts 600, 602 pulse-echo scanning 783–5 pulsed double electron–electron resonance (PELDOR) 560–1, 567–8, 575, 579 pulsed ESR techniques 560 pultrusion 138–9 pyrolysis 929–32, 945–6 pyrrolidine derivatives 553 pyrroline derivatives 553–4

1003

quantitative analysis of surfaces by electron spectroscopy (QUASES) 595–7 quantum dots 965 quasielastic behavior 741–2 quasielastic scattering 708, 710–13 R-cyclodextrin macrocycles (R-CD) 146–7 radiative heat transfer 388 radiochemical ageing 797, 809, 815, 818–21, 839 radiography 785–6 RAFT see reversible addition–fragmentation chain transfer Raman spectroscopy 922 random chain scission 808–12 random copolymers 261–3, 567, 628–9, 647 random phase approximation (RPA) 32, 50–2 rapid heat cooling (RHC) DSC 365 RAPRA standard scheme 819 Ratcliffe model 412–13 Rayleigh model 408 Rayleigh scattering 640 RBS see Rutherford backscattering spectrometry reaction-induced phase separation 649–51 reaction injection molding (RIM) 140–1, 157 reactive compatibilization 95–8, 197–8, 261 reactive coupling 94, 95–6 reactive diffusion 813 reactive flame retardancy 845 reactive grafts 600 reactive surfactants 96 reactor neutron sources 716–18 rear face temperature 391 recrystallization 90, 361, 373, 685–6 rectifying contacts 469–70 recycled thermoplastics 935–40, 949–51 recycling 10 feedstock 928–9, 945–6 mechanical 252, 296, 932–43, 947–51 polymer blends 11, 123 waste management 921–2, 924, 932–43, 945–51 red phosphorus 851, 855, 858 REDOR see rotational echo double resonance reflectometers 720 reinforced RIM (RRIM) 140–1 reinforcement see fibers; fillers relative hygrometry (RH) 828 relative permeability 754–7, 760–2 relaxation modulus curves 301–2 reptation chain dynamics 38 see also sticky reptation model resin dissolution 924–7

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resin transfer molding (RTM) 136–8, 157, 867–8, 889 resistance measurements 458–72 contact resistance 469–72 four-probes method 460–5 spreading resistance of contacts 468–70 two-probes method 459 Van der Pauw method 465–7 responsive interphases 101–5, 112 reversible addition–fragmentation chain transfer (RAFT) polymerization 99–100, 600 reversible adsorption theories 15–16, 27 reversible hydrolysis 828–9 RH see relative hygrometry RHC see rapid heat cooling rheological properties 9, 13–29, 311–57 ageing processes 810–13, 815–16, 827 blending laws 326–35 block copolymers 312–13, 339–54 capillary extrusion rheometry 347–53 concentrated suspensions 327–9 dilute suspensions 327 droplet breakup and deformation 318–22, 338 droplet coalescence 321–2, 323, 338 emulsions 330–5 entangled systems 14, 21–7 entanglement plateaux 340–1 experimental results 19–21, 25–7 flow-induced morphological changes 345–7 glass transition temperature 313–14 hydrodynamic modeling 13–14, 18 low frequency viscoelastic behavior 335–9 macroscopic characteristics 18, 24–5 microscopic characteristics 15–17, 22–4 miscible liquids 326 modeling and simulations 13–29, 326–39 morphological characteristics 311–18, 323–5, 334–40 polymer blends 311–18, 322–39 reinforcement mechanisms 14 shear and stress 317–18 shear-induced segregation 322–3 specific mechanical relaxation 343–5 specificity of blend rheology 322–6 sticky reptation model 22–5, 27 strain non-uniformity 322–3 suspensions of rigid particles 327, 333–5 swelling 325–6 thermorheological complexity 341–3 unentangled systems 14–21 uniform description of flow curves 329–30 viscoelasticity 316, 317, 326–39

viscosity ratio 315–17, 319–21, 322 waste management 933–6, 951 rheology-light scattering 644–5 rigid polymers 972 RIM see reaction injection molding robotic in-mold fiber reinforcement (RIMFIRE) 872 Rocking curves 734–5, 738–9 rope-like nanotubes 445–6 rotational echo double resonance (REDOR) 525 rotational molding 140–1 Rouse chain dynamics 38 RRIM see reinforced RIM RTM see resin transfer molding rubber blends/composites characterization 499–507, 657, 660–1 elastomers 964–5 electrical properties 456–7, 458 mechanical properties 253–5, 257–9, 265–6, 269–70, 276–9 morphological characteristics 185–6, 195, 201, 214 waste management 922, 924, 932, 937–8 rubber modified thermoplastics 9 rupture point 969 Rutherford backscattering spectrometry (RBS) 111 S-glass 880 Saito’s theory 808–9 SALS see small angle light scattering SANS see small angle neutron scattering SAXS see small angle X-ray scattering SBA-15 562–3 SBSS see short beam shear stresses scaling behavior 36–8 scanning electron microscopy (SEM) elastomers 963–4 electrical properties 432–3, 437–41, 451 manufacturing techniques 130, 133, 142–3, 146, 151, 155 mechanical properties 255, 257, 280, 300 morphological characteristics 8, 165 neutron scattering 733 rheological properties 334 waste management 939 X-ray scattering 672, 681–3, 693–6 scanning near-field optical microscopy (SNOM) 109–10 scanning thermal microscopy (SThM) 376 scanning tunnel microscopy (STM) 8 SCBA see self-contained breathing apparatus SCF see self-consistent field Schottky barrier diodes 910–11 SCL see shell crosslinked; space-charge-limited

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Index SCMF see single chain in mean field SCORIM see shear controlled orientation in injection molding scrap rubber tires (SRT) 130 SD see spinodal decomposition SDM see shear-induced demixing SEC see size exclusion chromatography secant moduli 260 secondary ion mass spectrometry (SIMS) 109, 610–11, 619–21, 627 secondary structure 523 segmental mobility 84, 92–3, 108–9 selective dissolution 314–15 selective multiphase materials 749, 767–9 self-assembly 39, 561–4 self-avoiding chains 36–7 self-concentration model 481 self-consistent field (SCF) theory 7, 31–2, 43–9 analytical theories 52, 55 fluctuation effects 48–9 mean field approaches 47 real systems 45–9, 68 theoretical application 43–5 self-contained breathing apparatus (SCBA) 876 self-correlation functions 712 self-similar grid structure theory 17 SEM see scanning electron microscopy semiconducting polymer composites 450–1, 458, 899 semi-empirical prediction models 412–15 semi-interpenetrating polymer networks (SIPN) 3–4, 570, 571–3 semi-IPNs 154–5 separate dispersions 193 separation of waste components 924–7 sequential elimination reactions 807, 818 sequential IPNs 151–2 series model 269–71, 275–6, 288–9, 404–5, 413–14 SFFT see single fiber fragmentation test shear controlled orientation in injection molding (SCORIM) 695–7 shear flow 168–72, 218–21, 651–5, 779 shear rates 70, 322–3 shear stress 317–18, 453 shear thinning/dynamics 730–3, 739–43 shear viscosity 19–20 shear-induced demixing (SDM) 651–2 shear-induced phase mixing (SIM) 651–2 shear-induced segregation 322–3 sheet molding compounds (SMC) 922, 924, 929–31, 944–51 sheeting formation mechanism 166–8, 172–4, 224–5

1005

shell crosslinked (SCL) micelles 629 short beam shear stresses (SBSS) 296 shrinkage 806, 830, 906 silane coupling agents 96, 98, 102 silica elastomers 962, 970 silica fumes 935 silica nanoparticles 26, 500–7 silicate layers see polymer-layered silicate nanocomposites silicification 967 silicon carbide 889 silicon containing polymers 891–3 silicone fillers 838 silk composites 136 silver fillers 437 silver ion determination 588, 600 SIM see shear-induced phase mixing SIMS see secondary ion mass spectrometry simulations see modeling and simulations simultaneous interpenetrating networks (SIN) 152–3 single chain in mean field (SCMF) simulations 64 single-chain partition functions 45–6 single fiber fragmentation test (SFFT) 88 single pulse excitation (SPE) MAS 531 single screw extruders (SSE) 129–31, 162, 164–5, 167, 184–5, 223–31 single vacuum bag (SVB) process 135–6 single-wall carbon nanotubes (SWNT) 227 applications 910 elastomers 966 electrical properties 444, 446–7 electron spin resonance spectroscopy 561–2 fire retardancy 859 SIPP see surface-initiated photopolymerization SIST see stepwise isothermal segregation technique site percolation 427 size exclusion chromatography (SEC) 156 skin–core interface 136, 816 SLS see static light scattering small angle light scattering (SALS) 640, 642–3, 646–57, 662 small angle neutron scattering (SANS) 8, 660, 670, 713–14, 719–20, 724–33 small angle X-ray scattering (SAXS) 670, 672–4, 677–87, 693–5, 697–700 manufacturing techniques 146 morphological characteristics 8 rheological properties 345 smart reversibly adaptive interphases 103–5 Snell’s law 715 SNOM see scanning near-field optical microscopy

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sodium bis(2-ethylhexyl) sulfosuccinate (AOT) 613–15 sol-gel process 145, 659–60, 960 solid-state NMR spectroscopy 110, 519–49 carbohydrates 530 chemical shifts 525–6 conducting polymers 530–1 copolymers 527–30, 542–3, 545–6 cross-polarization 519–20, 521–2, 525–36, 537 dipolar decoupling 521–2, 525, 536–7 drug delivery systems 535, 543–4 high-resolution 13 C NMR 521, 523–4, 525–36 magic angle spinning 519, 521–2, 525–36 morphological characteristics 537 NMR relaxation studies 538–44 other nuclei 536–8 polyesters 526–30 polymer blends 531–5, 539–42 polymer composites/nanocomposites 538, 543 polymer conformation 522–5 polymer–LMW compound interactions 535–6 polyolefins 526 polypeptides 536 solid polymers 525–36 spin diffusion 544–6 theoretical background 520–2 solid-state shear pulverization (SSSP) 162 solubility coefficients 749, 751 soluble-insoluble (SIS) polymers 104 solution blending 124–5 solution casting 257–8 solution intercalation 141–2 solvent absorption 802–3 solvolysis 928–9, 946–7 sorting of waste materials 922–4 space-charge-limited (SCL) emission 910 space charges 835–6 space intersection 792 spacecraft applications 891–4 spallation neutron sources 718 SPE see single pulse excitation specific heat capacity 9 definition 389 measurement techniques 392–3 polymer blends 396–400 polymer composites 395–404 temperature-dependence 395–401 unfilled polymer systems 395–400 specific mechanical relaxation 343–5 spin-coated humidity sensors 911–12 spin diffusion 544–6 spin probe/spin label techniques 552, 557, 563–5, 574

spin trapping agents 552 spinodal decomposition (SD) 40, 647–8 sporting goods 875–6 spray layup process 134–5 spreading resistance of contacts 468–70 spring-bead chain model 36, 59 SRIM see structural RIM SSA see successive self-nucleation and annealing SSSP see solid-state shear pulverization SST see strong stretching theory stabilizers 102, 933–4 starch/cellulose acetate (SCA) 902 starch/ethylene vinyl alcohol (SEVA) 902 starch/poly(acrylic acid) blends 544 static light scattering (SLS) 639 statistical percolation models 428–30 step-strain experiments 168, 170 stepwise breakup mechanism 318–20 stepwise crystallization 361, 364 stepwise isothermal segregation technique (SIST) 365 SThM see scanning thermal microscopy sticky reptation model 7, 22–5, 27 stimuli-responsive polymers (SRP) 104–5 STM see scanning tunnel microscopy storage moduli interpenetrating polymer networks 284–5, 289–90 polymer blends 269, 273–5, 336, 338, 341, 346–8 rheological modeling 20–1, 25–7 strain non-uniformity 322–3 strain-at-failure 258, 301 strain-induced crystallization 259, 262–3 strength-to-weight ratios 883–4 stress-induced crystallization 322 stress relaxation function 24 stress–strain behavior 9 elastomers 962, 964, 969 interpenetrating polymer networks 285 polymer blends 261–3, 268 polymer gels 292–5 X-ray scattering 694–5, 697–9 strong segregation limit (SSL) 561 strong segregation theory 32, 41, 46, 52–5 strong stretching theory (SST) 53–5 structural models 58–62 structural reorganization 800–2, 829–31 structural RIM (SRIM) 140–1 structure-oriented percolation models 430 structure–property relationships 479–80, 640 styrene-acrylonitrile copolymer (SAN) 95, 202–3, 211, 214, 565, 621–3, 647–8, 661

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Index styrene-butadiene rubber (SBR) elastomers 964–5 electron spin resonance spectroscopy 559, 574 manufacturing techniques 126–7, 128 mechanical properties 259, 265–6, 270 morphological characteristics 167, 209 styrene-butadiene-styrene (SBS) block copolymers 132, 254–6, 279, 312, 339–40, 347–50, 542 styrene copolymerization 146 styrene crosslinked polyesters (UP) 807 styrene-ethylene/butylene-styrene triblock (SEBS) copolymers morphological characteristics 185–6, 198–9, 205–6, 211, 228–9 rheological properties 313, 339–53 waste management 937 styrene-isobutylene-styrene (SIBS) 535, 544 successive self-nucleation and annealing (SSA) 365 sulfonation reactions 806–7 sulphur vulcanized polyisoprene (SVPIP) 799, 817 supercapacitors 910 supercritical carbon dioxide extraction 927 supercritical fluids 132–3 superficial degradation 815–16 surface discharges 836–7 surface energies 107–8 surface erosion 165–6, 172 surface-initiated photopolymerization (SIPP) 603 surface modified carbon nanotubes (CNT) 446–9 surface properties 67, 434, 592–7, 716–18 surface specificity of XPS 587 suspensions 327–9, 333–5 swelling processes 325–6, 961–3 synthetic metals see inherently conducting polymers systematic bottom-up modeling 57–8 tactoid formation 180–1, 220 Taffy process 609 Takayanaki model 288–9 targeted reactive interphase modifiers 96 Taylor theory 168–70, 187–8, 316–17, 319, 331 TCL see transcrystalline layers TDGL see time-dependent Ginzburg-Landau TDR see time domain reflectometry TDS see time domain spectroscopy tear properties 9 TEM see transmission electron microscopy temperature modulated DSC (TMDSC) 373–7 temperature-modulated TGA (MTGA) 378 temperature programmed desorption (TPD) 610 temperature-sensitive microgels 145

1007

tensile strength 9 interpenetrating polymer networks 284–6, 288–9 polymer blends 254–5, 260–4, 269–70 polymer gels 292–4 polymer nanocomposites 298–9 waste management 929–31, 934–6, 939–40 tetrabromobisphenol A (TBBPA) 843, 850, 927 tetrabromophthalic anhydride (TBPA) 850 tetraethoxysilane (TEOS) 537, 541, 960 tetramethoxy orthosilane (TMOS) 562–3, 659–60 tetramethyl bisphenol A polycarbonate (TMPC) 656 textile industry applications 866 TGA see thermal gravimetric analysis thermal ageing 797, 818–19, 821–3, 838 thermal analysis 359–85 differential scanning calorimetry 360, 362–77 evolved gas analysis 378–80 localized thermomechanical analysis 376–7 micro- and nanothermal analysis 376–7 polarized light microscopy 359–60 polymer blends 370–3 polymer composites 375 temperature modulated DSC 373–7 thermal gravimetric analysis 378–80 thermal optical microscopy 359–64 thermal conductivity 9 definition 388–9 measurement techniques 389–90, 393 modeling and simulations 404–19 polymer blends 396–400, 404 polymer composites 395–419 temperature-dependence 395–401 unfilled polymer systems 395–400 thermal degradation 531, 533, 834 thermal diffusivity definition 389 measurement techniques 390–2, 393 polymer blends 396–400 polymer composites 395–404 temperature-dependence 395–401 unfilled polymer systems 395–400 thermal effusivity 389, 393–4 thermal gravimetric analysis (TGA) 151, 378–80, 625–6 thermal interfacial interactions 85 thermal mapping 109 thermal optical microscopy (TOM) 359–64 thermal waste management processes 929–32, 945–6 thermally stimulated conductivity (TSC) 154 thermally stimulated depolarization current (TSDC) analysis 88, 154, 481–2, 485–9, 493–507, 512–13 thermography 786–8

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thermomechanical analysis (TMA) 154, 376–7 thermo-optical properties 892–3 thermophysical properties 1–5, 387–423 definitions 388–9 fillers 406–19 future developments 419–20 macroscopic 8–9 measurement techniques 389–94 modeling and simulations 395, 404–19 polymer blends 396–400, 419 polymer composites 387–8, 390, 394–420 simultaneous measurement of parameters 393–4 unfilled polymer systems 394–400, 419 thermoplastic elastomers (TPE) 312 thermoplastic IPNs 154 thermoplastic olefin (TPO) nanocomposites 216–17, 942–3 thermoplastics applications 868–9, 877, 879–80, 886, 907 characterization 9, 615–16, 670 life cycling 10 recycled 935–40, 949–51 waste management 924, 928, 932–43, 949–51 thermorheological complexity 341–3 thermoset adhesives 886–7 thermoset IPNs 282–3 thermosets ageing processes 812 applications 895, 907 nanofillers 972–3 thermal analysis 375 waste management 924, 927–8, 944–51 Thiele moduli 765 thiodipropionates 825–6 third harmonic (3ω) method 390 thread breakup 318 three-dimensional (3D) numerical modeling 417–19 three-point parameters 411 THS see transient hot strips THW see transient hot wire time-dependent Ginzburg-Landau (TDGL) theories 65 time domain reflectometry (TDR) 483 time domain spectroscopy (TDS) 483 time-of-flight (ToF) spectrometers 610–11, 619–21, 627, 720–1, 741 time-resolved dynamic light scattering (TR-DLS) 659–60 time resolved-SALS (TR-SALS) 640, 650–1, 656 tip-streaming 172, 202 tip-stretching 202, 224–5 tissue engineering 105, 898–900 titanium dioxide 853–4, 859, 961–2

TMA see thermomechanical analysis TMDSC see temperature modulated DSC ToF see time-of-flight TOM see thermal optical microscopy Tomotika model 321 topological networks 146–7 Torquato model 410–12 torque ratio see viscosity ratio tortuosity factors 754, 755–8, 761–2 Tougaard theory 595–6 toughness 810 TPD see temperature programmed desorption TR-DLS see time-resolved dynamic light scattering TR-SALS see time resolved-SALS train-loop-tail structures 82 transcrystalline interlayers 90–1, 111–12 transcrystalline layers (TCL) 361, 363–4, 684, 692–3, 831 transient hot strips (THS) 394 transient hot wire (THW) 394 transient plane source (TPS) 393–4 transmission electron microscopy (TEM) dielectric spectroscopy 500 elastomers 961, 963–4 electrical properties 446 manufacturing techniques 142–3, 145–6 mechanical properties 266–7, 281–3, 286 morphological characteristics 8, 187 rheological properties 334, 339, 345–7 waste management 942–3 transport properties 10 transporter interphases 101–2 triaryl phosphates 851, 861 triboelectric separation 922–3 tricalcium phosphate (TCP) 903–4 triethyleneglycol methacrylate (TREGMA) 145 triethylenetetramine (TETA) 609, 611–12, 626 trimethylene carbonate (TMC) 564 trimethylolpropane trimethacrylate (TMPTM) 565 triphenyl phosphate (TPP) 861 triple axis spectrometers 721–3 Tsao model 407 TSC see thermally stimulated conductivity TSDC see thermally stimulated depolarization current tungsten oxide 912 turbine blades 877–8 twin-screw extruders (TSE) 131–2, 164–7, 181–2, 186–7, 207–8, 220–31 two-dimensional (2D) numerical modeling 415–17 two-dimensional NMR spectroscopy 522 two-parameter Edwards model 36

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Index two-probes method 459 two-roll mixing 126–7 UL 94V flammability test 848, 853, 855, 860–1 ultra small-angle X-ray scattering (USAXS) 674 ultrasonic scanning 783–5 ultrathin polymer coatings 600–1 ultraviolet (UV) absorption spectroscopy 807–8 ultraviolet (UV) resistance 877 ultraviolet (UV) spectroscopy 8 unentangled systems 14–21 uniformly swollen gels 5 unsaturated polyesters (UP) 137, 139, 283, 826 upper critical solution temperature (UCST) 94, 646 vacuum bag molding 135–6, 157 vacuum-assisted resin infusion molding (VARIM) 877–8 vacuum-assisted resin transfer molding (VARTM) 136 Van der Pauw method 465–7 VE resins 283 vinyl polymers 399–400 Vinyloop process 924–5 viscoelastic emulsion model 333–5 viscoelasticity mechanical properties 251, 269, 272 modeling and simulations 7 morphological characteristics 171, 174–6, 187–92, 219 rheological modeling 13–14, 18, 22, 25–7 rheological properties 316, 317, 326–39 viscosity block copolymers 349–53 morphological characteristics 171, 175, 183–7, 214, 219 polymer blends 323, 328–32 rheological modeling 13, 19–20 waste management 933 viscosity ratio 315–17, 319–21, 322, 331 viscosity-shear rate 9 visual inspection 781–2 Vogel-Tammann-Fulcher (VTF) equation 484, 490 voltage stabilizers 837–8 volume fractions 331, 755, 962 vorticity elongation/breakup 172 vulcanized elastomers 812 waste hierarchy 922 waste management 921–57 feedstock recycling 928–9, 945–6 identification and sorting 922–4 interfacial agents 932–5 International context 944–5

1009

mechanical recycling 932–43, 947–51 polymer blends 932, 936–7 polymer composites 921–57 polymer nanocomposites 940–3 separation of components 924–7 thermal processes 929–32, 945–6 thermoplastics 924, 928, 932–43, 949–51 thermosets 924, 927–8, 944–51 water-induced damage 826–9, 832–3, 836–7, 891 water trees 835 water vapor permeability 752 wavelength dispersive spectrometers (WDX) 587 weak segregation limit 41, 46, 50–2, 55–6 weight average molar mass 810–12, 828–9 wetting processes 452 wide angle X-ray scattering (WAXS) 8, 670, 672–7, 680–2, 687–97, 699–700 wide-angle X-ray diffraction (WAXD) 124, 187 Wiener measure 35 Wilhelmy-plates 108 wind turbine blades 877–8 wobbling 190 Wollaston probes 376 wollastonite–silver (W-Ag) 408–9, 439 wormlike chain model 35–6 wormlike micelles 282, 283 wrapping 449 X-radiography 785–6 X-ray diffraction (XRD) 142, 155 X-ray fluorescence correlation spectroscopy (XFCS) 110 X-ray photoelectron spectroscopy (XPS) 109, 585–637 applications to polymeric materials 599–630 binding energy values 589, 591 chemical shifts 588–91 colloidal particles 605–8 conducting polymers 613–19 copolymers 589–90, 618–19, 628–9 data acquisition and processing 599 elemental analysis 588 epoxy resins 609–12, 627 grafted polymers 589–90, 600–5 instrumentation 597–9 interpenetrating polymer networks 627–8 manufacturing techniques 142 morphological characteristics 8 overlayer thickness 592–7 photoionization 586–7 polymer blends 619–24 polymer composites/nanocomposites 624–7 principles of technique 586–99

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Index

X-ray photoelectron spectroscopy (Continued) quantification 592 spectral examination and analysis 587–91 spectrometer design 597–8 surface specificity 587 X-ray scattering (SAXS/WAXS) 669–703 crystallinity of polymers/polymer systems 669–70, 673–8, 695 microfibrillar reinforced composites 670, 671–2, 678, 681–93 morphological characteristics 8 polymer blends 670, 693–5 polymer composites/nanocomposites 670, 671–3, 678–93, 695–9 rheological properties 345 SCORIM molded nanocomposites 695–7 stress-strain curves 694–5, 697–9 techniques and principles 669–71 theoretical background 671–8 xerogels 145–6

yield strength 9 yield stress 259, 281, 801–2 Young’s moduli ageing processes 801, 830, 835 interpenetrating polymer networks 284, 288–9 polymer blends 256, 260–5, 271 waste management 930, 936, 951 YP compatibilizers 681–93

Zeeman interactions 555 zeolites 90–2, 767–9, 967 zero injection pressure resin transfer molding (ZIP RTM) 872 zeta potentials 108 Zimm chain dynamics 38 zinc borates 850, 855 zip elimination reactions 807, 818 zirconia 961