Rubber-Clay Nanocomposites: Science, Technology, and Applications

RUBBER-CLAY NANOCOMPOSITES

RUBBER-CLAY NANOCOMPOSITES Science, Technology, and Applications Edited by MAURIZIO GALIMBERTI

Copyright Ó 2011 by John Wiley & Sons, Inc. All rights reserved. Published by John Wiley & Sons, Inc., Hoboken, New Jersey. Published simultaneously in Canada. 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, scanning, or otherwise, except as permitted under Sections 107 or 108 of the 1976 United States Copyright Act, without either the prior written permission of the Publisher, or authorization through payment of the appropriate per-copy fee to the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, (978) 750-8400, fax (978) 750-4470, or on the web at www.copyright.com. Requests to the Publisher for permission should be addressed to the Permissions Department, John Wiley & Sons, Inc., 111 River Street, Hoboken, NJ 07030, (201) 748-6011, fax (201) 748-6008, or online at http://www.wiley.com/go/permission. Limit of Liability/Disclaimer of Warranty: While the publisher and author have used their best efforts in preparing this book, they make no representations or warranties with respect to the accuracy or completeness of the contents of this book and specifically disclaim any implied warranties of merchantability or fitness for a particular purpose. No warranty may be created or extended by sales representatives or written sales materials. The advice and strategies contained herein may not be suitable for your situation. You should consult with a professional where appropriate. Neither the publisher nor author shall be liable for any loss of profit or any other commercial damages, including but not limited to special, incidental, consequential, or other damages. For general information on our other products and services or for technical support, please contact our Customer Care Department within the United States at (800) 762-2974, outside the United States at (317) 572-3993 or fax (317) 572-4002. Wiley also publishes its books in a variety of electronic formats. Some content that appears in print may not be available in electronic formats. For more information about Wiley products, visit our web site at www.wiley.com. Library of Congress Cataloging-in-Publication Data: Rubber-clay nanocomposites : science, technology, and applications / edited by Maurizio Galimberti. p. cm. Includes bibliographical references and index. ISBN 978-0-470-56210-9 1. Nanocomposites (Materials) 2. Rubber. I. Galimberti, Maurizio. TA455.R8R825 2011 620.1’94–dc23 2011021008 Printed in the United States of America oBooK ISBN: 978-1-118-09286-6 ePDF ISBN: 978-1-118-09288-0 ePub ISBN: 978-1-118-09287-3

CONTENTS PREFACE

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CONTRIBUTORS

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SECTION I 1

CLAYS FOR NANOCOMPOSITES

CLAYS AND CLAY MINERALS 1.1 1.2

1.3

1.4

3

What’s in a Name / 3 Multiscale Organization of Clay Minerals / 6 1.2.1 Dispersion Versus Aggregation / 6 1.2.2 Delamination/Exfoliation Versus Stacking / 6 Intimate Organization of the Layer / 8 1.3.1 Cationic and Neutral Clay Minerals / 8 1.3.2 Anionic Clay Minerals (O) / 21 Most Relevant Physicochemical Properties of Clay Mineral / 22 1.4.1 Surface Area and Porosity / 22 1.4.2 Chemical Landscape of the Clay Surfaces / 24 1.4.3 Cation (and Anion) Exchange Capacity / 24 1.4.4 Intercalation and Confinement in the Interlayer Space / 27 1.4.5 Swelling / 30 1.4.6 Rheology / 31 v

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1.5 Availability of Natural Clays and Synthetic Clay Minerals / 33 1.6 Clays and (Modified) Clay Minerals as Fillers / 35 Acknowledgment / 37 References / 37 2

ORGANOPHILIC CLAY MINERALS

45

2.1

Organophilicity/Lipophilicity and the Hydrophilic/Lipophilic Balance (HLB) / 45 2.2 From Clays to Organoclays in Polymer Technology / 47 2.3 Methods of Organoclay Synthesis / 49 2.3.1 Cation Exchange from Solutions / 49 2.3.2 Solid-State Intercalation / 58 2.3.3 Grafting from Solution / 59 2.3.4 Direct Synthesis of Grafted Organoclays / 62 2.3.5 Postsynthesis Modifications of Organoclays: The “PCH” / 64 2.3.6 An Overview of Commercial Organoclays / 64 2.3.7 One-Pot CPN Formation / 66 2.4 Other Types of Clay Modifications for Clay-Based Nanomaterials / 66 2.4.1 Organo-Pillared Clays / 66 2.4.2 Plasma-Treated Clays / 69 2.5 Fine-Tuning of Organoclays Properties / 69 2.5.1 Maximizing the Dispersion of the Filler: Effect of Surfactant/CEC Ratio / 69 2.5.2 Improving Thermal Stability / 70 2.5.3 Chemical Treatments / 71 2.5.4 Physical Treatments (Freeze-Drying, Sonication, Microwave) / 71 2.6 Some Introductory Reflections on Organoclay Polymer Nanocomposites / 72 References / 75 3

INDUSTRIAL TREATMENTS AND MODIFICATION OF CLAY MINERALS 3.1

3.2

Bentonite: From Mine to Plant / 87 3.1.1 A Largely Diffused Clay / 87 3.1.2 Geological Occurrence / 89 3.1.3 Mining / 89 Processing of Bentonite / 90 3.2.1 Modification of Bentonite Properties / 90 3.2.2 Processing Technologies / 91

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3.3

Purification of Clay / 93 3.3.1 Influence of Clay Concentration / 94 3.3.2 Influence of Swelling Time / 94 3.3.3 Influence of Temperature / 95 3.4 Reaction of Clay with Organic Substances / 97 3.5 Particle Size Modification / 99 References / 99 4

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

101

4.1

Structure and Dynamics / 101 4.1.1 Packing Density and Self-Assembly / 102 4.1.2 Dynamics and Diffusion at the Clay–Surfactant Interface / 110 4.1.3 Utility of Molecular Simulation to Obtain Molecular-Level Insight / 111 4.2 Thermal Properties / 111 4.2.1 Reversible Melting Transitions of Alkyl Chains in the Interlayer / 111 4.2.2 Solvent Evaporation and Thermal Elimination of Alkyl Surfactants / 113 4.3 Layer Separation and Miscibility with Polymers / 115 4.3.1 Thermodynamics Model for Exfoliation in Polymer Matrices / 115 4.3.2 Cleavage Energy / 116 4.3.3 Surface Energy / 121 4.4 Mechanical Properties of Clay Minerals / 121 References / 123 5

CHEMISTRY OF RUBBER–ORGANOCLAY NANOCOMPOSITES 5.1 5.2

Introduction / 127 Organic Cation Decomposition in Salts, Organoclays and Polymer Nanocomposites / 128 5.2.1 Experimental Techniques / 128 5.2.2 Decomposition of Organoclays Versus Precursor Organic Cation Salts / 133 5.3 Mechanism of Thermal Decomposition of Organoclays / 135 5.4 Role of Organic Cations in Organoclays as Rubber Vulcanization Activators / 137 References / 141

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SECTION II

6

PREPARATION AND CHARACTERIZATION OF RUBBER–CLAY NANOCOMPOSITES

PROCESSING METHODS FOR THE PREPARATION OF RUBBER–CLAY NANOCOMPOSITES

147

6.1 6.2

Introduction / 147 Latex Compounding Method / 148 6.2.1 Mechanism / 148 6.2.2 Influencing Factors / 149 6.3 Melt Compounding / 157 6.3.1 Mechanism / 157 6.3.2 Influencing Factors / 160 6.4 Solution Intercalation and In Situ Polymerization Intercalation / 170 6.5 Summary and Prospect / 170 Acknowledgment / 171 References / 171 7

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES 7.1

7.2

7.3

7.4

Introduction / 181 7.1.1 Focus, Objective and Structure of Chapter 7 / 181 7.1.2 X-Ray Diffraction Analysis for the Investigation of RCN / 182 Background for the Review of RCN Morphology / 182 7.2.1 Cationic Clays Used for the Preparation of Rubber Nanocomposites / 182 7.2.2 Multiscale Organization of Layered Clays / 184 7.2.3 Clay Distribution and Dispersion / 184 7.2.4 Clay Modification: Intercalation of Low Molecular Mass Substances / 184 7.2.5 Types of Polymer–Clay Composites / 184 7.2.6 Specific Literature on RCN / 186 Rubber–Clay Nanocomposites with Pristine Clays / 186 7.3.1 Rubber Nanocomposites with Cationic Clays / 187 7.3.2 In a Nutshell / 187 7.3.3 Distribution and Dispersion of a Pristine Clay in a Rubber Matrix / 190 7.3.4 Organization of Aggregated Pristine Clays / 194 Rubber–Clay Nanocomposites with Clays Modified with Primary Alkenylamines / 197 7.4.1 In a Nutshell / 197 7.4.2 Composites with Montmorillonite and Bentonite / 198

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7.4.3

Composites with Fluorohectorite Modified with a Primary Alkenylamine / 202 7.5 Rubber–Clay Nanocomposites with Clays Modified with an Ammonium Cation Having three Methyls and One Long-Chain Alkenyl Substituents / 206 7.5.1 In a Nutshell / 206 7.5.2 Composites with Montmorillonite and Bentonite / 207 7.6 Rubber–Clay Nanocomposites with Montmorillonite Modified with Two Substituents Larger Than Methyl / 212 7.6.1 In a Nutshell / 212 7.6.2 Hydrogenated Tallow and Benzyl Groups as Ammonium Cation Substituents / 213 7.6.3 Hydrogenated Tallow and Ethylhexyl Groups as Ammonium Cation Substituents / 213 7.6.4 Other Long- and Short-Chain Alkenyl Groups as Ammonium Cation Substituents / 215 7.7 Rubber Composites with Montmorillonite Modified with an Ammonium Cation Containing a Polar Group / 215 7.7.1 In a Nutshell / 217 7.7.2 Composites with Diene Rubbers / 217 7.8 Rubber Nanocomposites with Montmorillonite Modified with an Ammonium Cation Containing Two Long-Chain Alkenyl Substituents / 219 7.8.1 In a Nutshell / 220 7.8.2 Composites with Two Talloyl Groups as Ammonium Cation Substituents / 220 7.9 Proposed Mechanisms for the Formation of Rubber–Clay Nanocomposites / 228 7.9.1 Two Mechanisms for the Formation of an Exfoliated Clay / 228 7.9.2 Two Mechanisms for the Formation of an Intercalated Organoclay / 228 7.9.3 Intercalation of Polymer Chains in the Interlayer Space / 229 7.9.4 Intercalation of Low Molecular Mass Substances in the Interlayer Space / 230 Abbreviations / 232 Acknowledgment / 233 References / 233 8

RHEOLOGY OF RUBBER–CLAY NANOCOMPOSITES 8.1 8.2

Introduction / 241 Rheological Behavior of Rubber–Clay Nanocomposites / 242

241

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8.2.1

Natural Rubber (NR), Epoxidized Natural Rubber (ENR) and Polyisoprene Rubber (IR)–Clay Nanocomposites / 243 8.2.2 Styrene–Butadiene Rubber (SBR)–Clay Nanocomposites / 246 8.2.3 Polybutadiene Rubber (BR)–Clay Nanocomposites / 247 8.2.4 Acrylonitrile Butadiene Rubber (NBR)–Clay Nanocomposites / 250 8.2.5 Ethylene Propylene Rubber–Clay Nanocomposites / 253 8.2.6 Fluoroelastomer–Clay Nanocomposites / 254 8.2.7 Poly(isobutylene-co-para-methylstyrene) (BIMS) Rubber–Clay Nanocomposites / 257 8.2.8 Poly(ethylene-co-vinylacetate) (EVA) Rubber–Clay Nanocomposites / 257 8.2.9 Polyepichlorohydrin Rubber–Clay Nanocomposites / 259 8.2.10 Thermoplastic Polyurethane (TPU)–Clay Nanocomposites / 261 8.2.11 Styrene–Ethylene–Butylene–Styrene (SEBS) Block Copolymer–Clay Nanocomposites / 262 8.3 General Remarks on Rheology of Rubber–Clay Nanocomposites / 263 8.4 Overview of Rheological Theories of Polymer–Clay Nanocomposites / 269 8.5 Conclusion and Outlook / 270 References / 271 9

VULCANIZATION CHARACTERISTICS AND CURING KINETIC OF RUBBER–ORGANOCLAY NANOCOMPOSITES 9.1 9.2 9.3

Introduction / 275 Vulcanization Reaction / 276 Rubber Cross-Linking Systems / 278 9.3.1 Sulfur Vulcanization / 278 9.3.2 Peroxide Vulcanization / 282 9.4 The Role of Organoclay on Vulcanization Reaction / 283 9.4.1 Influence of Organoclay Structural Characteristics on Rubber Vulcanization / 288 9.5 Vulcanization Kinetics of Rubber–Organoclay Nanocomposites / 290 9.6 Conclusions / 297 References / 298

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10

MECHANICAL AND FRACTURE MECHANICS PROPERTIES OF RUBBER COMPOSITIONS WITH REINFORCING COMPONENTS

xi

305

10.1 10.2

Introduction / 305 Testing of Viscoelastic and Mechanical Properties of Reinforced Elastomeric Materials / 307 10.2.1 Dynamic–Mechanical Analysis / 307 10.2.2 Tensile Testing / 310 10.2.3 Assessment of Toughness Behavior under Impact-Like Loading Conditions / 313 10.2.4 Hardness Testing / 315 10.2.5 Special Methods / 316 10.3 Characterization of the Fracture Behavior of Elastomers / 319 10.3.1 Fracture Mechanics Concepts / 319 10.3.2 Experimental Methods / 321 10.4 Mechanism of Reinforcement in Rubber–Clay Composites / 328 10.5 Theories and Modeling of Reinforcement / 333 Acknowledgment / 336 References / 336 11

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY 11.1

Introduction / 343 11.1.1 Butyl Rubbers as Nanocomposite Base Elastomers / 343 11.1.2 Measurement of Tire Innerliner Compound Permeability / 345 11.1.3 Further Improvement in Tire Permeability / 346 11.2 Nanocomposites / 346 11.3 Preparation of Elastomer Nanocomposites / 352 11.4 Temperature and Compound Permeability / 352 11.5 Vulcanization of Nanocomposite Compounds and Permeability / 356 11.6 Thermodynamics and BIMSM Montmorillonite Nanocomposites / 358 11.7 Nanocomposites and Tire Performance / 362 11.8 Summary / 364 References / 364

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SECTION III 12

COMPOUNDS WITH RUBBER–CLAY NANOCOMPOSITES

RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

369

12.1 12.2

Introduction / 369 Preparation Methods / 371 12.2.1 Latex / 371 12.2.2 Solution / 373 12.2.3 Melt Blending / 374 12.3 Cure Characteristics / 377 12.4 Clay Dispersion / 379 12.4.1 Detection / 380 12.4.2 Characterization / 383 12.5 Properties / 387 12.5.1 Mechanical (Dynamic–Mechanical) / 387 12.5.2 Friction/Wear/Abrasion / 392 12.5.3 Barrier / 393 12.5.4 Fire Resistance / 396 12.5.5 Others / 397 12.6 Applications and Future Trends / 398 Acknowledgment / 399 References / 399 13

RUBBER–CLAY NANOCOMPOSITES BASED ON NITRILE RUBBER 13.1 13.2

Introduction / 409 Preparation Methods and Clay Dispersion / 410 13.2.1 Solution / 410 13.2.2 Latex / 411 13.2.3 Melt Blending / 412 13.3 Cure Characteristics / 414 13.4 Properties / 416 13.4.1 Mechanical (Dynamic–Mechanical) / 416 13.4.2 Friction/Wear / 421 13.4.3 Barrier / 423 13.4.4 Fire Resistance / 424 13.4.5 Others / 425 13.5 Outlook / 425 Acknowledgment / 426 References / 426

409

14

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RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

431

14.1

Introduction / 431 14.1.1 Butyl Rubber: Key Properties and Applications / 431 14.1.2 Butyl Rubber–Clay Nanocomposites / 433 14.2 Types of Clays Useful in Butyl Rubber–Clay Nanocomposites / 435 14.2.1 Montmorillonite Clays / 435 14.2.2 Hydrotalcite Clays / 435 14.2.3 High Aspect Ratio Talc Fillers / 436 14.2.4 Other Clays / 437 14.3 Compatibilizer Systems for Butyl Rubber–Clay Nanocomposites / 438 14.3.1 Surfactants and Swelling Agents / 439 14.3.2 Butyl Rubber Ionomers / 439 14.3.3 Maleic Anhydride-Grafted Polymers / 443 14.3.4 Low Molecular Weight Polymers and Resins / 444 14.4 Methods of Preparation of Butyl Rubber–Clay Nanocomposites / 444 14.4.1 Melt Method / 445 14.4.2 Solution Method / 445 14.4.3 Latex Method / 447 14.4.4 In Situ Polymerization / 448 14.5 Properties and Applications of Butyl Rubber–Clay Nanocomposites / 449 14.5.1 Air Barrier Properties / 449 14.5.2 Reinforcement Properties / 452 14.5.3 Vulcanization Properties / 454 14.5.4 Adhesion Properties / 456 14.5.5 Other Properties / 457 14.6 Conclusions / 457 References / 458 15

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM) 15.1 15.2 15.3 15.4

Introduction / 465 Types of Clay Minerals Useful in EPM–, EPDM–Clay Nanocomposites / 466 Compatibilizer Systems for Olefinic Rubber–Clay Nanocomposites / 467 Preparation of EPDM–Clay Nanocomposites by an In Situ Intercalation Method / 469

465

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15.5

Characteristics of EPDM–Clay Nanocomposites / 473 15.5.1 Gas Barrier Properties of EPDM–Clay Nanocomposites / 473 15.5.2 Rheological Properties of EPDM–Clay Nanocomposites / 474 15.5.3 Stability of EPDM–Clay Nanocomposites / 475 15.5.4 Swelling Properties of EPDM–Clay Nanocomposites / 475 15.5.5 Mechanical Properties of EPDM–Clay Nanocomposites / 476 15.6 Preparation and Characteristics of EPM–Clay Nanocomposites / 479 15.6.1 Tensile Properties of EPM–CNs / 480 15.6.2 Temperature Dependence of Dynamic Storage Moduli of EPM–CNs / 481 15.6.3 Creep Properties of EPM–CNs / 482 15.6.4 Swelling Properties of EPM–CNs / 483 15.7 Conclusions / 486 References / 486 16

RUBBER–CLAY NANOCOMPOSITES BASED ON THERMOPLASTIC ELASTOMERS 16.1 16.2

16.3

16.4 16.5

Introduction / 489 Selection of Materials / 491 16.2.1 Polymer Resin / 491 16.2.2 Nanoparticles / 493 Experimental / 493 16.3.1 Processing of Thermoplastic Elastomer Nanocomposites / 493 16.3.2 Morphological Characterization / 494 16.3.3 Thermal Properties Characterization / 495 16.3.4 Flammability Properties Characterization / 495 16.3.5 Thermophysical Properties Characterization / 496 Numerical / 497 16.4.1 Modeling of Decomposition Kinetics / 497 Discussion of Results / 501 16.5.1 Nanoparticle Dispersion / 501 16.5.2 Thermal Properties / 503 16.5.3 Flammability Properties / 507 16.5.4 Microstructures of Posttest Specimens / 511 16.5.5 Thermophysical Properties / 512 16.5.6 Kinetic Parameters / 513

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16.6 Summary and Conclusions / 516 16.7 Nomenclature / 517 Acknowledgments / 518 References / 518

SECTION IV 17

APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES

AUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES 17.1 17.2

Introduction / 525 Automotive Application of Rubber / 526 17.2.1 Automotive Hose / 527 17.2.2 Automotive Seals / 528 17.2.3 Automotive Belts / 529 17.2.4 Automotive Tubing / 529 17.2.5 Door Seal and Window Channels / 529 17.2.6 Diaphragms and Rubber Boots / 529 17.2.7 Tire, Tube and Flap / 529 17.2.8 Other Miscellaneous Rubber Parts / 531 17.3 Prime Requirement of Different Elastomeric Auto Components from Application Point of View / 531 17.4 Elastomeric Nanocomposites and Rubber Industry / 531 17.5 Superiority of Clay/Clay Mineral in Comparison to Other Nanofillers / 534 17.6 Organo-Modified Clay/Clay Minerals / 534 17.7 Scope of Application of Elastomeric Nanocomposites in Automotive Industry / 534 17.7.1 Lighter Weight and Balanced Mechanical Property / 535 17.7.2 Barrier Property or Air Retention Property / 538 17.7.3 Aging and Ozone Resistance / 539 17.7.4 Solvent Resistance / 541 17.7.5 Better Processability / 542 17.7.6 Elastomeric Polyurethane–Organoclay Nanocomposites / 544 17.7.7 Use of Organoclay Nanocomposites in Tire / 545 17.8 Disadvantages of Use of Organoclay Elastomeric Nanocomposites in Automotive Industry / 548 17.9 Conclusion / 549 Acknowledgment / 550 References / 550

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18

CONTENTS

NONAUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES

557

18.1

Water-Based Nanocomposites / 557 18.1.1 Barrier Properties / 557 18.1.2 Comparison with Thermally Processed Elastomers / 566 18.2 Applications / 566 18.2.1 Sports Balls and Other Pneumatic Applications / 566 18.2.2 Breakthrough Time Applications / 571 References / 573 INDEX

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PREFACE

A NEW GENERATION OF FILLERS FOR RUBBERS: NANO-FILLERS The properties of rubbers have always fascinated the human mind. “I wonder why the night, as a rubber, is of endless elasticity and softness” wrote a novelist [1] and the inventor of vulcanization, Charles Goodyear, reported “There is probably no other inert substance the properties of which excite in the human mind an equal amount of curiosity, surprise, and admiration. Who can examine and reflect upon this property of gum-elastic without adoring the wisdom of the Creator?” [2]. Rubbers are indeed fundamental materials for the human life but their properties are not sufficient for their applications, not even after vulcanization. To achieve the required physical–mechanical properties, rubbers have to be reinforced with the so-called reinforcing fillers. The addition of carbon black was observed to improve the physical properties of vulcanized rubbers already at the beginning of twentieth century, although its use was delayed by consumers resistance to the black color. The first synthetic rubber tires in a car, of Emperor Wilhelm, were of a white color. Fillers such as carbon black and silica were developed all over the last century and allowed to improve a large set of rubber properties, such as impermeability, tear, fatigue and abrasion resistance, while simultaneously increasing antagonistic properties such as modulus and elongation at break: this phenomenon is known as the paradox of elastomers. Carbon black and silica are made of spherical primary particles, with an average size in a range from 5 to 100 nm, and are present in the rubber matrix as aggregates that

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cannot be separated via thermomechanical mixing, having dimensions up to several hundreds of nanometers. Over the past two decades, new characters appeared on the scene of fillers for polymeric materials, the so-called nano-fillers: they can be dispersed in a polymer matrix as individual particles with at least one dimension at the nanoscale. This nanometric size is correlated with features such as a huge specific surface area, a very low concentration for establishing a network in a polymer matrix (what is known as the percolation threshold) and also, often, a high length-to-width ratio, that is, a high aspect ratio. Most researches performed both in the academic and industrial fields on polymer composites with nano-fillers, that is, on polymer nano-composites aimed to exploit the enormous potential of nano-fillers. CLAYS AS NANO-FILLERS FOR RUBBERS Among nano-fillers, clays undoubtedly play a major role. These layered silicates are available as inexpensive natural minerals and, as in the case of the most diffused cationic clay, montmorillonite, are considered to have a safe toxicologic profile, as they appear to have little chance to cross biological barriers. They are thus suitable for large scale applications and were used to prepare novel rubber/inorganic materials. Clays are hydrophilic and need to be compatibilized with the hydrocarbon rubber matrix: the most applied organophilic modifiers, ammonium cations bearing long-chain alkenyl substituents, are able to build up a variety of crystalline arrangements in the interlayer space. The so-formed organically modified clays promote a multiscale organization in the rubber matrix, from its distribution and dispersion to a reorganization of the organic moiety between two opposite layers, potentially involving the polymer chains. The onium modifiers are also known as efficient accelerators of the cross-linking reactions. Moreover, all these aspects depend, to a different extent, on the type of rubber adopted as the matrix. All these aspects are degrees of freedom but represent at the same time a complexity for the development of these novel rubber materials. Rubber–clay nanocomposites (RCN) have been extensively investigated. Hundreds of papers are available in the scientific literature, with a large number of data and some proposed interpretations. Some industrial applications have already been successfully brought to a commercial scale, backed by hundreds of patent applications and based in particular on the improvement of mechanical properties as well as of impermeability. However, in spite of the large scientific investigation and of some commercial applications, the potential of clays in imparting new properties to a rubber composite could be exploited to a much larger extent. WHY A BOOK ON RUBBER–CLAY NANO-COMPOSITES? This book moves from the awareness of the state of the art of RCN and its identity is determined by the following objectives. To make available an updated recollection of

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data, interpretations and theories reported in the open scientific literature. Time is mature for proposing a rationalization of what so far discovered, to the benefit of both students and professionals. A further objective is to allow scientists and technologists working in the field to critically review the common perception of RCN, building a sound cultural base, prodrome of further R&D activities and further innovations. As a key feature of this book, items involved in RCN science, technology, and applications are discussed providing a comprehensive overview from clay structural features to application (e.g., in an automotive part). This book wishes thus to contribute to a better exploitation of RCN potential. A BRIEF SUMMARY OF THE BOOK This book is organized in four sections. In the first section, clays and organoclays for rubber composites are introduced. In Chapter 1, Bergaya et al. present natural and synthetic clay minerals, from crystallographic structure to fundamental aspects such as the multiscale clay organization and, in particular, the intimate organization of the layers. Most relevant clay physicochemical properties are also discussed. Clay modification with the preparation of organoclays is covered by the same authors in Chapter 2, analyzing the finetuning of organoclays properties. The industrial treatments of a bentonite clay is discussed by Della Porta in Chapter 3: processing, purification, reaction with organic substances. Heinz illustrates in Chapter 4, the alkylammonium chains on layered clay mineral surfaces: structure and dynamics, thermal and mechanical properties, layer separation, and miscibility with polymers. Giannini et al. deal with chemistry of rubber–organoclay nanocomposites in Chapter 5, from their thermal decomposition to the interaction with the sulfur-based vulcanization chemistry. The second section is dedicated to preparation and characterization of RCN. Zhang et al. present in Chapter 6 the processing methods for the preparation of RCN, in particular latex and melt compounding, from mechanism to influencing factors. Galimberti et al. rationalize the RCN morphology in Chapter 7: the multiscale organization in the rubber matrix is discussed for pristine clays and organoclays, as a function of processing method, type of rubber, and in particular of the organic modifier. Mechanisms proposed for the formation of intercalated and exfoliated clays are critically reviewed. Isaev et al. deal with RCN rheology in Chapter 8, taking into consideration various types of rubbers and providing an overview of proposed theories. Vulcanization characteristics and curing kinetic of RCN are discussed by Lopez Manchado et al. in Chapter 9, focusing the attention on the role of organoclay in a vulcanization reaction and on the influence of its structural characteristics. The mechanical and fracture mechanics properties of RCN are reviewed by Reincke et al. in Chapter 10, dealing with viscoelastic and mechanical properties, fracture behavior and mechanisms, theories and modeling of reinforcement. The permeability of RCN is covered by Rodgers et al. in Chapter 11, with particular reference to butyl type rubbers, to influencing factors such as rubber vulcanization and temperature and to an important application such as the one in a tire compound.

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The third section is dedicated to RCN based on a particular type of rubber. Karger-Kocsis et al. deal with apolar diene rubber and with nitrile rubber in Chapters 12 and 13, respectively. RCN based on butyl and halobutyl rubbers are covered by Magill et al. in Chapter 14. Makoto and Koo et al. discuss RCN based on olefinic rubbers and thermoplastic elastomers in Chapters 15 and 16, respectively. Preparation methods are covered, key aspects such as barrier, vulcanization, mechanical properties are discussed. In the final section, main applications of RCN are presented. Bandyopadhyay et al. discuss automotive applications in Chapter 17 and Feeney et al. present in Chapter 18 nonautomotive applications such as the one for sport balls. Last but not the least, as the editor I wish to acknowledge the work of all the authors, done with much involvement and enthusiasm. We felt as a team, with the common aim to give a profitable contribution to all the readers. I would like finally to aknowledge the work done by my coworkers, Valeria Cipolletti and Michele Coombs, in editing this book. REFERENCES 1. Yoshimoto, B. Asleep, Grove/Atlantic, Inc, New York, 2000. 2. Goodyear, C. Gum Elastic and its Variation with a Detailed Account of its Applications and Uses, New Haven, 1855, Vol. 1.

CONTRIBUTORS DANA ADKINSON, Lanxess, Inc., Butyl Rubber Global Research and Development, London, Ontario, Canada SAMAR BANDYOPADHYAY, Hari Shankar Singhania Elastomer and Tyre Research Institute, Rajsamand, Rajasthan, India FAI¨ZA BERGAYA, CRMD-CNRS, University of Orleans, Orleans, France NATACHA BITINIS, Instituto de Ciencia y Tecnologıa de Polımeros, CSIC, Madrid, Spain MORGAN C. BRUNS, Department of Mechanical Engineering, The University of Texas at Austin, Austin, Texas, USA SUGATA CHAKRABORTY, Hari Shankar Singhania Elastomer and Tyre Research Institute, Rajsamand, Rajasthan, India JAESUN CHOI, Institute of Polymer Engineering, University of Akron, Akron, Ohio, USA VALERIA ROSARIA CIPOLLETTI, Pirelli Tyre S.p.A., Milan, Italy ATTILIO CITTERIO, Dipartimento di Chimica, Materiali e Ingegneria Chimica, Politecnico di Milano, Milan, Italy DAFNE COZZI, Dipartimento di Chimica, Materiali e Ingegneria Chimica, Politecnico di Milano, Milan, Italy CINZIA DELLA PORTA, Laviosa Chimica Mineraria S.p.A., Livorno, Italy xxi

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CONTRIBUTORS

OFODIKE. A. EZEKOYE, Department of Mechanical Engineering, The University of Texas at Austin, Austin, Texas, USA CARRIE FEENEY, InMat Inc., Hillsborough, New Jersey, USA MAURIZIO GALIMBERTI, Pirelli Tyre S.p.A., Milan, Italy; and Dipartimento di Chimica, Materiali e Ingegneria Chimica, Politecnico di Milano, Milan, Italy KONSTANTINOS G. GATOS, Megaplast S.A., Research & Development Center, Athens, Greece LUCA GIANNINI, Pirelli Tyre S.p.A., Milan, Italy SIMONA GIUDICE, Dipartimento di Chimica, Materiali e Ingegneria Chimica, Politecnico di Milano, Milan, Italy HARRIS A. GOLDBERG, InMat Inc., Hillsborough, New Jersey, USA WOLFGANG GRELLMANN, Center of Engineering Sciences, Martin Luther University of Halle-Wittenberg, Halle, Germany HENDRIK HEINZ, Department of Polymer Engineering, University of Akron, Akron, Ohio, USA MARIANELLA HERNA´NDEZ-SANTANA, Instituto de Ciencia y Tecnologıa de Polımeros, CSIC, Madrid, Spain WAI K. HO, Department of Mechanical Engineering, The University of Texas at Austin, Austin, Texas, USA AVRAAM I. ISAYEV, Institute of Polymer Engineering, University of Akron, Akron, Ohio, USA MAGUY JABER, LRS-CNRS, University of Paris, Paris, France JO´ZSEF KARGER-KOCSIS, Tshwane University of Technology, Pretoria, South Africa; and Budapest University of Technology and Economics, Budapest, Hungary MAKOTO KATO, Toyota Central R&D Labs, Inc., Nagakute, Aichi, Japan JOSE` MARIA KENNY, Instituto de Ciencia y Tecnologıa de Polımeros, CSIC, Madrid, Spain JOSEPH H. KOO, Department of Mechanical Engineering, The University of Texas at Austin, Austin, Texas, USA JEAN-FRANC¸OIS LAMBERT, LRS-CNRS, University of Paris, Paris, France JASON C. LEE, Department of Mechanical Engineering, The University of Texas at Austin, Austin, Texas, USA DAVID J. LOHSE, ExxonMobil Research & Engineering Co. 1545 Route 22 East P. O. Box 998 Annandale, NJ 08801-3059

CONTRIBUTORS

xxiii

MIGUEL ANGEL LOPEZ-MANCHADO, Instituto de Ciencia y Tecnologıa de Polımeros, CSIC, Madrid, Spain YONG-LAI LU, Beijing University of Chemical Technology, Beijing, China CHARLES PHILIPPE MAGILL, Lanxess, Inc., Butyl Rubber Global Research and Development, London, Ontario, Canada RABINDRA MUKHOPADHYAY, Hari Shankar Singhania Elastomer and Tyre Research Institute, Rajsamand, Rajasthan, India KATRIN REINCKE, Center of Engineering Sciences, Martin Luther University of Halle-Wittenberg, Halle, Germany BRENDAN RODGERS, ExxonMobil Chemical Company, Baytown, Texas, USA RALF I. SCHENKEL, Lanxess, Inc., Butyl Rubber Global Research and Development, London, Ontario, Canada JOHN SOISSON, ExxonMobil Chemical Company, Baytown, Texas, USA RAQUEL VERDEJO, Instituto de Ciencia y Tecnologıa de Polımeros, CSIC, Madrid, Spain WALTER WADDELL, ExxonMobil Chemical Company, Baytown, Texas, USA ROBERT WEBB, ExxonMobil Chemical Company, Baytown, Texas, USA WEIQING WENG, ExxonMobil Chemical Company, Baytown, Texas, USA LI-QUN ZHANG, Beijing University of Chemical Technology, Beijing, China

SECTION I

CLAYS FOR NANOCOMPOSITES

CHAPTER 1

CLAYS AND CLAY MINERALS FAÏZA BERGAYA MAGUY JABER JEAN-FRANÇOIS LAMBERT

1.1 WHAT'S IN A NAME The term “clay” was used in everyday language long before being imbued with a well-defined scientific meaning. Therefore, it is not surprising that it carries different connotations to different communities. To the industrialist, it is a raw material available in large amounts at cheap prices, characterized by its macroscopic properties relative to various applications. To the geologist working in the field, it is a particular secondary mineral largely found in weathered deposits from sedimentary or volcanic origin. To the chemist and mineralogist, it refers to a particular type of mineral structure defined at the atomic level. Recent recommendations of the JNC1 advise to use the term “clay minerals” to refer to precisely determined crystallographic structures, and define “clays” in terms of macroscopic properties.2 Therefore, a natural clay will consist of a/several clay minerals mixed with additional minerals as impurities. However, the distinction is not always clearly made and many papers that use well-defined clay minerals will refer to them as “clays” because the full denomination is somewhat cumbersome [1]. Here we will build on the crystallographic view, which is the most rigorous, and try to indicate how the atomic structure dictates the properties at other levels. The most salient structural feature of clay minerals is that they are layered. That is to say they belong to a large class of inorganic compounds built by the stacking of 1

Joint Nomenclature Committee of the AIPEA (Association Internationale pour l’Etude des Argiles) and the CMS (Clay Minerals Society). 2 A clay is “a naturally occurring material composed primarily of fine-grained minerals, which is generally plastic at appropriate water contents and will harden when dried or fired.”

Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

3

CLAYS AND CLAY MINERALS

4

Iono-covalent interactions

Basal distance

(b)

(a) H-bonds or dispersion interactions

(c)

Interlayer

– + –

– + –

– + –

+ – +

+ – +

+ – +

FIGURE 1.1 The basic architecture of a clay mineral at the nanometric scale: (a) neutral layers; (b) negatively charged layers with compensating cations (cationic clays); and (c) positively charged layers with compensating anions (anionic clays).

two-dimensional units, known as layers, whose internal coherence is due to strong iono-covalent bonds, while in the direction perpendicular to the stacking they are bound to each other through weaker forces. This means that the layers can be separated from each other relatively easily, and the volume included between two successive layers, whatever its content, is called the “interlayer space” (the term “gallery” or “intergallery space” was formerly used synonymously but has to be discarded). A macroscopic analogy would be a pile of paper sheets, or may be a deck of cards depending on the semirigidity assumed for the layers. Clay minerals are a subset of the family of layered oxides (or oxyhydroxides), which can be classified in three different categories according to the electrical charge of the layer (Figure 1.1): (i) Neutral layers, as in pyrophyllite, talc, and kaolinite. The layers are held together by van der Waals interactions and/or hydrogen bonds. (ii) Negatively charged layers. Since the structure as a whole must be neutral, the negative layer charge must be compensated exactly by an equal amount of positive charges provided by cations located in the interlayer space (compensating cations). These minerals are most often listed as phyllosilicates, and the most widespread in nature (especially montmorillonite which is the major component of commercial bentonites) belong to this group, and are therefore called “cationic clays” when specification is needed. (iii) Positively charged layers with compensating anions in the interlayer space. The most common natural mineral in this group is hydrotalcite (HT), but this belongs to a broader family of “HT-like” materials most often synthesized in the laboratory and called layered double hydroxides or LDH. These are also called “anionic clays” [2–4]. It should be noted that we have characterized clay minerals as layered “oxides” or “oxyhydroxides” rather than layered silicates (or “phyllosilicates”). Indeed some of them do not contain any silicon in their formula (the LDH) and thus are certainly not silicates. Even in the case of cationic clays, an argument could be made that the term “silicate” obscures the real structure of the layers (as outlined in Section 1.3.1.3 and

WHAT'S IN A NAME

5

corresponding inset) and is a leftover from a time when only the raw formula was known. Let us come now to the clay as a macroscopic material. Historically, the criterion of particle size has been used a lot to define clays, although different disciplines and professions have fixed different size limits. The “clay fraction” has been defined as fine-grained materials with a maximum particle size (or, rather, an equivalent spherical diameter)
2 mm. However, the particle size limit used by different communities could vary from 1 mm (for colloid scientists) up to 4 mm (for engineers; see Ref. [5]. It is not considered as good practice any more to set a well-defined size limit to clay minerals [1], although the particles must be small enough to form colloidal dispersions in water. More than a definition, the size criterion is a practical recipe for separation since particles with different sizes will sediment at different rates in water, according to Stokes’s law [6] (which states that the settling velocity of a particle in a fluid is proportional to the square of the particle radius, all other things being equal) — the finest clay fraction can thus be obtained by centrifugation at the laboratory scale, even though industrialists use other methods to separate large quantities of clays for practical reasons (see Chapter 3). Materials commercially available as clays may be (i) raw clays, containing several other associated minerals (carbonates, cristobalite, feldspars, quartz, etc.) and other associated X-ray amorphous phases (organic matter, iron hydroxides, etc.) as contaminants in addition to clay minerals proper or (ii) clay mineral fractions obtained by sedimentation, fragmentation, and/or several other treatments aimed, for example, at eliminating iron oxide impurities by selective dissolution. The recommended procedures to obtain pure clay mineral samples are reported in more details by Carrado et al. [7]. For some applications, de novo synthesis of clay minerals may be more attractive compared to the complexity of purification processes of the raw clays. Synthetic cationic clay minerals may indeed be prepared with high purity (cf. Section 1.5), but their cost will of course be higher than that of natural clays. The JNC has kept among the clay-defining criteria some referring to specific properties that are relevant to their application as materials: plasticity in the presence of water, hardening on drying. In these properties, the behavior of clays in the presence of water plays an important role. We will therefore consider the behavior of clays in the presence of aqueous phases, first at a nanoscopic level, which includes mesoscopic level as shown in Figure 1.2, and later at a molecular level. “Nanoscale” (1–100 nm) Molecular range

1Å

Mesoscopic range (2–50 nm)

1 nm

100 nm

Size range (logarithmic scale) FIGURE 1.2 Mesoscale, nanoscale, and molecular scale.

6

CLAYS AND CLAY MINERALS

1.2 MULTISCALE ORGANIZATION OF CLAY MINERALS Clay minerals form pastes, then gels, when the amount of water is increased or other polar solvents are added to the dry solid, and that translates a specific organization of the dispersed solids in the presence of the solvent. That organization is largely similar to the one that exists in native clays; in the simplest scheme, three successive levels of organization can be defined at different scales: in descending order of size, aggregates, particles, and layers [1,8]. 1.2.1 Dispersion Versus Aggregation At the upper level (macroscopic), the sample is made of millimetric-size aggregates. Upon closer inspection, the aggregates are seen to consist of a number of flat micrometer size particles (sometimes called platelets; the formerly used term of “tactoids” is rather ill-defined and should be avoided). Here “aggregation” is used in opposition to dispersion to mean the agglomeration of clay particles that results at the macroscopic level in the visual observation of flocculation (or coagulation; both terms are used indifferently); flocculation may be followed by sedimentation, or not. In water dispersion, particles and even single layers can be associated in different ways depending on solution conditions, especially on the pH value and ionic strength: face-to-face (FF, the most frequent arrangement), edge-to-edge (EE), or edge-to-face (EF). Extended EF agglomerates are sometimes called “house-ofcards” or cardhouse structures (Figure 1.3). This level of organization is important to understand the swelling (cf. Section 1.4.5) and rheological (cf. Section 1.4.6) properties of clays. The aspect ratio is defined as the average ratio of the width to the thickness of the particles; values of 5–30 are typical, although for completely delaminated samples the aspect ratio might exceed 100. Many factors can induce a strong tendency of clay particles to aggregate: high clay concentration, high ionic strength (concentration of ionized salts), presence of surfactants and organic polymers. The nature of the compensating cations plays a particularly important role in the aggregation/dispersion behavior: strongly hydrated cations, that is, cations with small radius such as Liþ , induce the dispersion of the aggregates into small nanoparticles. Finer structural details such as the charge heterogeneity (cf. Section 1.4.2) of the clay surfaces, or the redox state of iron (Fe2þ or Fe3þ ) also play an important role on dispersion/flocculation. Sometimes it is desirable to increase the dispersion of a clay/water system. Vigorous mechanical agitation of a dispersion containing low clay concentration is needed. Several techniques are used involving industrial devices as extruders, internal mixers, ultrasonicators, and so on (see Chapter 3). 1.2.2 Delamination/Exfoliation Versus Stacking At the lowest (molecular) level, the individual particles shown in Figure 1.3 are composed of the stacking of elementary layers, alternating with interlayer spaces

MULTISCALE ORGANIZATION OF CLAY MINERALS

(a)

(b)

1 µm

1 µm

7

(c)

FIGURE 1.3 Top: aggregation of clay particles at the micrometer scale, mostly edge-to-edge (EE) (a) the inset shows the stacking of elementary layers within a particle; and mostly edge-toface (EF) (b). Bottom: higher level organization of aggregates (c).

containing the compensating cations, and whatever other molecules may happen to be intercalated (cf. Section 1.4.4). The stacking exhibits crystallographic periodicity along the c axis, even though successive layers may be oriented differently according to the a and b directions. Depending on the environmental conditions, their size varies from a few stacked layers (2–5) to much greater numbers. Some conditions favor a separation of the individual layers (cf. Section 1.4.4), the term of delamination is used to designate the separation between the planar faces of two adjacent layers. The layers may eventually become completely independent from one another, with a loss of crystallographic orientation; each unit is then freely oriented in space, independently from the others. We propose to call this particular stage exfoliation, although no clear distinction is made between delamination and exfoliation in many papers. Exfoliation occurs when the delaminated units (including isolated layers or stackings of a few layers) are isotropically dispersed in the aqueous or solvent matrix. This is observed for instance in aqueous dilute Laponite dispersions. Figure 1.4 shows an electron micrograph of an exfoliated Laponite dispersion where the macroscopic heterogeneity of the sample is apparent (exfoliated layers coexist with stackings of a few layers). The state of stacking/delamination depend not only on the considered clay mineral and on the dispersion medium but also on the thermodynamic conditions (pressure and temperature, pH, and ionic strength), just as was the case for the higher level of association, namely, aggregation (cf. Section 1.2.1).

8

CLAYS AND CLAY MINERALS

FIGURE 1.4 TEM micrograph of a dispersed Laponite.

In clay minerals of the smectite group, the more the layers are stacked, the thicker and more rigid the particles are. The thickest particles appear as “flakes” (e.g., for montmorillonite) or as “laths” (e.g., for hectorite), while completely exfoliated clay layers are flexible. The degree of stacking is, perhaps obviously, related to the specific surface area (SSA) of the clay mineral as measured by N2 physisorption (the latter is of course applied to dry samples, and not to samples in aqueous dispersions), since the interlayer spaces are not accessible to N2 sorption, which therefore takes place only on the external surface of the particles. For a typical montmorillonite, consisting of micrometric particles and stacking of a few tens of layers, one calculates that the developed surface area ranges between 10 and 20 m2/g. This indeed corresponds to the low range of values observed experimentally (structural defects, cracks, etc., may lead to higher values).

1.3 INTIMATE ORGANIZATION OF THE LAYER The broad definition of clay minerals presented above (Sections 1.1 and 1.2) does not provide information on the molecular structure of the units designated as layers, except that they must have a strong bidimensional cohesion. We will now consider in more detail some of the crystallographic structures that qualify as clay minerals. 1.3.1 Cationic and Neutral Clay Minerals An important class of lamellar compounds of natural or synthetic origin involves the cationic clay minerals belonging to the phyllosilicate group [9,10]. 1.3.1.1 General Organizational Principles The organization of phyllosilicates is particularly rich and well studied and we will briefly outline it here.

INTIMATE ORGANIZATION OF THE LAYER

9

Each layer of the structure is in fact constituted by the assembly of two or three sheets, which are either tetrahedral or octahedral. Tetrahedral sheets are abbreviated as “T,” and they are constituted of cornersharing [XO4] units, where X is a small cation, which may be either Si4þ or Al3þ , although other substitutions are possible; oxide ions (formally O2  ) occupy the corners. Octahedral sheets, abbreviated as “O,” consist of edge-sharing [MO4(OH)2] units, where M can be either a trivalent (such as Al3þ ), a divalent (such as Mg2þ ), or a monovalent (Liþ ) ion; the central site of the octahedron may also be vacant. The structures of tetrahedral and octahedral sheets, showing the connectivity of the elementary units, are shown in Figure 1.5. The central feature explaining clay architecture is due to the fact that the repeat distances of the hexagonally symmetric tetrahedral and octahedral sheets are almost exactly coincident, which allows the outward-pointing oxygen of each tetrahedron in the tetrahedral sheet to be shared with the octahedral sheet. A first classification can be proposed based on the layer type, that is, the particular succession of sheets building the layer: (i) In 1 : 1 or TO type clay minerals, the layer is formed by one tetrahedral sheet linked to one octahedral sheet. (ii) In 2 : 1 or TOT type clay minerals, two tetrahedral sheets are linked to both sides of a central octahedral sheet. This very frequent “sandwich” structure will be illustrated in Figure 1.6, which represents a montmorillonite (see below). Note that some phyllosilicates such as chlorites have a main TOT layer with the same structure as above, which alternates with another octahedral sheet (bruciteor gibbsite-like) in the interlayer. They are also considered as TOT type [9,10]. The use of “TOTO” or “2:1:1” to designate this layer type, often found in the ancient literature, should be discarded. The second criterion is the occupancy of the octahedral sheet leading to (i) clay minerals with “trioctahedral character”, where all the octahedral sites are occupied by a dication such as Mg2þ ; (ii) clay minerals with “dioctahedral character”, where 2/3 of octahedral sites are occupied by a trication such as Al3þ ion and the third octahedral site is empty. Note that “trioctahedral” clay minerals contain dications, and conversely. The third and last criterion is the charge per formula unit for each layer, which is due to isomorphic substitutions within the layers. Substitutions of high-charge cations by lower charge ones intrinsically generate a deficit of positive charges in the layers, which are therefore negatively charged. These negative charges must be counterbalanced for the structure as a whole to be electrically neutral. This is ensured by “compensating cations” in the interlayer.

10

CLAYS AND CLAY MINERALS

FIGURE 1.5 Showing the connectivity, in the tetrahedral sheet (a) of a clay layer and in the octahedral sheet (b) (for a dioctahedral mineral, where one of three octahedra is vacant).

INTIMATE ORGANIZATION OF THE LAYER

11

FIGURE 1.6 The structure of montmorillonite, showing two successive layers. The interlayer space is occupied by compensating cations with varying degrees of hydration.

The substitutions may be chiefly present in the octahedral sheet (e.g., Al3þ replaced by Mg2þ or Fe2þ , or Mg2þ replaced by Liþ ), or in the tetrahedral sheet (e.g., Si4þ replaced by Al3þ ). This source of variability will be explored in more detail below, in the case of the smectite group. Tetrahedral substitutions generate localized layer charges, while the charges generated by octahedral substitutions are smeared out by the octahedral sheets on both sides and may be considered as delocalized on the layer surface. 1.3.1.2 The Main Clay Minerals Groups (TO and TOT) Taken together, the three criteria given above lead to nine clay minerals main groups (detailed below). It must be underlined that this classification is based on the molecular-level order and does not allow predicting the larger scale morphologies (plates, laths, fibers, rings, or nanotubes). The raw formulas of some representative clay minerals are given in Table 1.1. Since they may be confusing for the nonspecialist, especially because different conventions may be used to write them, the question of clay formula interpretation is addressed in a separate inset (see p.16.) (I) The kaolinite and serpentine group, typical TO phyllosilicates, where the charge of the two-sheets layer is almost zero. The best-known species of the first subgroup are kaolinite (a planar phyllosilicate structure) and halloysite (a spheroidal phyllosilicate structure made of nanotubes). In the serpentine group, the best-known species is probably chrysotile (a rolled phyllosilicate structure). As indicated by the examples of halloysite and chrysotile,thetwo-dimensional connectivity of the layers does notguarantee that platelets will form at the upper organizational level; structural

12

0.6–0.9

0.2–0.6

0

0

Charge per Formula Unit (Half-Unit Cell)

þ

IV

VI

þ

ðSi4  x Alx Þ ðAl2  y Mgy ÞO10 ðOHÞ2 , (x þ y)K

VI

þ

IV. Vermiculite group

ðSi4  x Alx Þ ðAl2 ÞO10 ðOHÞ2 , x M  nH2 O

Vermiculite

IV

Beidellite

ðSi4 Þ ðAl2  y Mgy ÞO10 ðOHÞ2 , y M  nH2 O

VI

Montmorillonite

IV

ðSi4 Þ VI ðMg3  y Liy ÞO10 ðOHÞ2 , yM þ  nH2 O ðSi4  x Alx Þ VI ðMg3 ÞO10 ðOHÞ2 , xM þ  nH2 O

IV

ðSi4  x Alx Þ VI ðMg3  y M3y þ ÞO10 ðOHÞ2 , (x  y)/2 M þ

Vermiculite

IV

Saponite

IV

Hectorite

ðSi4 Þ VI ðMg3 ÞO10 ðOHÞ2

IV

ðSi4 Þ ðAl2 ÞO10 ðOHÞ2

Talc

IV

VI

Pyrophyllite

III. Smectite group

II. Talc–pyrophyllite group

ðSi2 Þ VI ðMg3 ÞO5 ðOHÞ4

IV

ðSi2 Þ ðAl2 ÞO5 ðOHÞ4

Serpentine

VI

IV

I. Serpentine–kaolin group

Trioctahedral Species

Kaolinite

Dioctahedral Species

TABLE 1.1 Layer Charge and Idealized Formula of Some Representative 1:1 and 2:1 Clay Minerals

13

ðSi2 Þ VI ðAl2 ÞO5 ðOHÞ4

Imogolite IV ðSiÞ VI ðAl2 ÞO3 ðOHÞ4

IV

Allophane

ðSi4  x Alx Þ VI ðMg3  y Liy ÞO10 ðOHÞ2 , (x þ y)K þ ðSi3 AlÞ VI ðMg3 ÞO10 ðOHÞ2 , K þ

IV

VIII. Allophane and imogolite group

Mg-palygorskite to Al sepiolite

IV

ðSi6 Þ VI ðMg4 yM2þ y ÞO15 ðOHÞ2 , 6H2O

Sepiolite

Tri–tri (trioctahedral layer and trioctahedral interlayer sheet)

ðSiAl3 Þ VI ðMg2 AlÞO10 ðOHÞ2 , Ca2 þ

Clintonite

IV

Phlogopite

IV

Lepidolite

Left-hand superscripts refer to cation location in the tetrahedral (IV) or in the octahedral (VI) layer.

Variable

Di (-tri) palygorskite

Variable

(Di–tri) and (tri–di)

VI. Chlorites group

þ

VII. Palygorskite and sepiolite group

ðSi4  x Alx Þ VI AlMgÞO10 ðOHÞ2 , 6H2O

Di–di (dioctahedral layer and dioctahedral interlayer sheet)

IV

þ

ðSi2 Al2 Þ VI ðAl2 ÞO10 ðOHÞ2 , Ca2 þ

Margarite

IV

VI

ðSi3 AlÞ ðAl2 Þ O10 ðOHÞ2 , K

VI

Muscovite

IV

VI

ðSi4  x Alx Þ ðFe2  y Mgy ÞO10 ðOHÞ2 , (x þ y)K

Celadonite

IV

Variable

1.8–2.0

0.9–1.0

V. True (flexible) and brittle mica group

14

CLAYS AND CLAY MINERALS

(II)

(III)

(IV)

(V)

(VI)

(VII)

constraints in the layers may cause them to curl and adopt different morphologies. The pyrophyllite and talc group are nonswelling TOT phyllosilicates without isomorphous substitutions; thus, the charge of the three-sheets layer is almost zero. The only species of this group are pyrophyllite (dioctahedral) and talc (trioctahedral), but the following groups can be understood as derived from them by increasing degrees of substitution. The group of smectites3 is composed of TOT planar phyllosilicates; they are also known as swelling clay minerals. Their name comes from the Greek “smhktikoV,” whose initial meaning is “cleaning earth.” The charge of the three-sheet layers varies from 0.2 to 0.6 per half unit cell. According to the importance of the smectite in the nanocomposites technology, only this group will be detailed further (Section 1.3.1.3). Figure 1.6 presents the structure of montmorillonite, a typical smectite. The vermiculite group, TOT phyllosilicates with a more limited swelling ability as compared to the smectites group and where the charge of the three-sheet layers varies from 0.6 to 0.9 per half unit cell. The most frequent vermiculite species are trioctahedral. The true (flexible) micas and brittle micas groups, TOT phyllosilicates where the charge of the three-sheet layers varies from 0.9 to 2 per half unit cell. The most common mica, that is, illite is subject of controversy in the literature (see comment on p. 15 in Ref. [1] and pp. 39–40 in Ref. [10]). Illite is considered either as a species of the true mica group (lying at the borderline with vermiculite as regards the degree of substitution) or as a separate group [11,12]. Illite is thought to be a mineral derived from smectite dehydration, progressing from montmorillonite to beidellite to illite [13]. This probably explains the occurrence in nature of interstratified smectite–illite layers presented below. The chlorites group, currently considered as TOT phyllosilicates (see Section 1.3.1.1 above). Here, the TOT layer bearing a net negative charge alternates with a single octahedral sheet bearing a positive charge in the interlayer space. Trioctahedral chlorites are the most common, where both the TOT and the interlayer sheet are trioctahedral, but other combinations also exist [10]. Interstratified Clay Minerals Group. The term of interstratified mineral is used to designate a lamellar material where different types of layers may be found in the stacking. It is a common phenomenon that may be considered as an intergrowth of different types of layers along the c axis (i.e., along the stacking direction). In fact, different combinations occur in nature between all the previous TO and/or TOT groups, in a regular or irregular manner, leading to different kinds of more complex layer types. For example,

3 Their morphology is two-dimensional, like that of “smectic” liquid crystals, whose name obviously comes from the same root.

INTIMATE ORGANIZATION OF THE LAYER

15

irregular illite–smectite (I–S) are often encountered in mineralogy, where the particles contain variable proportions of swelling and nonswelling interlayers [14,15]; one also finds kaolinite–smectite (K–S) systems that consist of an alternating irregular layer sequence of kaolinite (TO) layers with smectite (TOT) layers. These combinations are called irregular interstratified clay minerals. However, when the succession of different layers occurs in a regular manner, a specific name is generally attributed to these interstratified stacked sequence [15] as, for example, rectorite that consists of a regular ordered succession of dioctahedral smectite and illite layers and may be symbolized as I–S–I–S [16]. It should be noticed that another kind of interstratification occurs when the simple smectite particles contain water interlayers of different thicknesses (cf. Section 1.4.4.2). This source of interstratification (filling interstratification) should not be confused with the above-mentioned structural interstratification of mixed layers systems. (VIII) The sepiolite and palygorskite group with the TOT layer-fibrous structure. In opposition to the previously mentioned groups, this one presents only one continuous two-dimensional tetrahedral sheet but a discontinuous octahedral sheet. This structure contains in fact fragments of TOT structures that extend along the a axis (Figure 1.7). The two representative species are trioctahedral sepiolite and dioctahedral palygorskite, which differ by their unit cell dimension (larger in the case of sepiolite, with larger channels also than for palygorskite). However, a recent study has shown that intermediate compositions exist between these two species [17]. The term “attapulgite” that was given by de Lapparent [18] to a clay mineral discovered in fullers’ earth from Attapulgus in the United States, has sometimes been used as synonymous with palygorskite and is still largely

FIGURE 1.7 The structure of sepiolite. Note the channels perpendicular to the plane of representation, which are filled with water molecules under ambient conditions.

16

CLAYS AND CLAY MINERALS

FIGURE 1.8 The structure of imogolite.

used in industry. Since connections in the direction perpendicular to the layers are assured in part by covalent bonds, minerals of this group cannot present the phenomenon of swelling that is defined in Section 1.4.5. (IX) Allophane and Imogolite Group. These aluminosilicates belong to the TO group (at least regarding short-range order) and are frequently found together in soils derived from volcanic ash. Both are very poorly crystalline minerals, being X-ray amorphous, and have an interesting fibrous morphology where the layers curl up to form 3–5 nm rings (allophane) or 2 nm tubes (for imogolite, the tubes being a few micrometers in length with an internal diameter of 1 nm). Allophane should actually be considered as a group since its composition is highly variable. The structure of imogolite is represented in Figure 1.8. It has been originally proposed by Cradwick et al. in 1972 [19]. Recently, this group of minerals has attracted renewed interest as it has been shown that imogolite-like materials (containing germanium instead of silicium) can be synthesized in large quantities [20], while double-walled Al–Ge imogolite-like nanotubes have been described; it is hoped that they could constitute a cheaper alternative to carbon nanotubes in some applications [21]. One can broadly distinguish all the previously described groups using their basal distances (d001) as a criterion. The TO groups are characterized by a distance of about 0.7 nm as for kaolinite (however, halloysite that is a hydrated kaolinite includes one molecular layer water and thus has a d001 of 1.0 nm). Allophane and imogolite have rather broad X-ray diffraction (XRD) peaks since they are poorly crystalline materials. For example, allophane shows XRD peaks at about 0.3 nm. The TOT groups exhibit basal distances of 1, 1.2, and 1.4 nm for mica, smectite, and chlorite groups, respectively. In the case of smectites, the observed value of 1.2 nm corresponds to an average distance, since it depends on the hydration state of the interlayer cations, which in turn may vary according to the treatment to

INTIMATE ORGANIZATION OF THE LAYER

17

which the sample has been submitted (cf. Section 1.3.1.3). Thus, one has to be careful in interpreting X-ray diffractograms of a clay sample; a d001 peak at 1.0 nm does not necessarily indicate the presence of mica, but might correspond to a fully dehydrated smectite. For the sepiolite and palygorskite group the basal distance is 1.2 nm. For the interstratified groups the situation is more complex, except for regular interstratified layers where a high distance is observed depending on the constituent layers. 1.3.1.3 The Swelling Clays: Smectites–Montmorillonites–Bentonites In view of the importance of smectites in the polymer nanocomposites technology, and their interesting swelling properties, this group will be studied in somewhat more detail. It includes the clay minerals most commonly used as rubber additives so far: montmorillonites, saponites, and hectorites. The smectites group can be divided into dioctahedral and trioctahedral species as shown in Table 1.2.

TABLE 1.2 Dioctahedral and Trioctahedral Species Belonging to the Smectite Group with Their Idealized Formulaa Dioctahedral Series

Trioctahedral Series

Montmorillonitea (Mt) IV a

Hectoritea (Ht)

ðSi4 Þ ðAl2  y Mgy ÞO10 ðOHÞ2 ðMn þ Þy=n VI

Bentonite: commercial name of a raw clay containing at least 50% of Mt (or other smectite)

Beidellite (Bd) IV

VI

ðSi4  x Alx Þ ðAl2 ÞO10 ðOHÞ2 ðM

IV a

ðSi4 Þ VI ðMg3  y Liy ÞO10 ðOHÞ2 ðMn þ Þy=n

Laponite and Fluoro-Ht are synthetic Ht (in Fluoro-Ht, the OH  ions are replaced by F  )

Saponite (Sap) nþ

Þx=n

V

ðSi4  x Alx Þ VI ðMg3 ÞO10 ðOHÞ2 ðMn þ Þx=n

Other species depending on the octahedral cation, which are different from Al or Mg Nontronite IV

Stevensite VI

ðSi4  x Alx Þ ðFe2 ÞO10 ðOHÞ2 ðM

nþ

Þx=n

Volkonskoite IV ðSi4  x Alx Þ VI ðCr2 ÞO10 ðOHÞ2 ðMn þ Þx=n

IV

ðSi4 Þ VI ðMg3  y &y ÞO10 ðOHÞ2 ðMn þ Þy=n (& represents a vacancy)

Sauconite ðSi4  x Alx Þ VI ðZn3 ÞO10 ðOHÞ2 ðMn þ Þx=n

IV

a The given formulas are somewhat idealized because the (x) and (y) substitutions located in the octahedral and tetrahedral sheets, respectively, are not rigorously limited to a single type of site in reality. Thus, even for Mt that ideally has only octahedral substitution, a minor amount of tetrahedral substitution may exist as well, giving rise to further structural variability such as montmorillonite with more or less pronounced beidellitic character. This is discussed in Ref. [22] together with many finer points of smectite classification.

CLAYS AND CLAY MINERALS

18

INTERPRETING THE CHEMICAL FORMULA OF A CLAY MINERAL

The chemical formulas of clay minerals are complex and may give the unpleasant impression that “anything goes,” especially since different conventions are used in the literature. Let us take as an example the formula of a calcium montmorillonite: IV

ðSi4 Þ VI ðAl2  y Mgy ÞO10 ðOHÞ2 ðMnþ Þy=n  mH2 O

ðTable 1:2Þ

It must be understood that such a formula successively lists the following: .

. .

The structural cations belonging to the layer, starting with those in the tetrahedral sheets (4Si4þ ) followed by those in the octahedral sheet (two cations overall, mostly Al3þ ; the stoichiometric coefficient (2  y), means that there is an unspecified degree (y) of Al3þ substitution, in this case by Mg2þ ). Note that cations belonging to the same sheet are associated by the use of brackets. In addition, we have used left-superscript roman numerals (VI and IV) to clearly identify the coordinance (number of neighbors) of the cations, VI in the octahedra and IV in the tetrahedra. The anions, most often O2  and OH  ; F  may substitute for OH  . The compensating cations, located in the interlayer region, which must be in the right amount to ensure electrostatic neutrality; sometimes, they are associated with interlayer water, written at the end of the formula as mH2O.

Often the coordinance of the ions is not explicitly indicated and the different constituents may be listed in a different order, for example, with compensating cations at the beginning. It is then left for the reader to identify in which sheet each group of ions is located—this is easily done since Si, which is always present, is to be found exclusively in the tetrahedral sheet. Sometimes, right instead of left superscripts are used for the coordinance but this runs the risk of inducing unfortunate confusions with the accepted notation for the oxidation number; thus, one has to understand that AlVI denotes six-coordinated aluminum, not the nonexistent Al6þ ion! The above formula corresponds to the contents of half a unit cell, which is the most frequent choice. Sometimes the formula is represented per (full) unit cell, and the formula of our calcium montmorillonite could then be given as Si8 ðAl4  z Mgz ÞO20 ðOHÞ4 Cax=2  nH2 O ðwhere z ¼ 2y; n ¼ 2mÞ Both conventions make sense crystallographically, and there is no point in seeking the stoichiometrically most simple formula since it would not be very enlightening for the understanding of the structure. At any rate, the structure as a whole must be electrically neutral and this is easily checked in many instances since most ions found in clay minerals have a single stable oxidation number under standard conditions—with some exceptions, the most conspicuous being iron that can be present as Fe2þ or Fe3þ .

INTIMATE ORGANIZATION OF THE LAYER

19

Montmorillonite (Figure 1.9 and structure in Figure 1.6) is the most studied in literature and the most used in different applications. While one will often find it in the literature designated as Mtm, MMT, and so on, it is advised to abbreviate this mineral name as Mt. The best-known montmorillonite-based material in the world is the mineral exploited in Wyoming (USA) at Fort Benton. The raw clay has been given the trade name of “bentonite” by Knight in 1898 [23]; it was marketed in 1920 by Baroid Corp., which later became NL Industries. The term of bentonite was introduced in Europe 10 years later [24] and has then been extended over the world to designate all the raw clays containing at least 50% of smectite and particularly of Mt. In fact, bentonites from Wyoming are somewhat atypical in that they contain Naþ as compensating cations, while most other bentonites known in the world are saturated by Ca2þ , and this causes very peculiar rheological properties. Usually, these Ca2þ bentonites are ion exchanged by sodium salts to convert them into their sodic form. This is called “sodium activated” bentonite.4 However, some properties such as viscosity remain different from the natural sodium Wyoming bentonite. For more information on Bentonites, see Refs [25] and [26]. It should also be mentioned that organobentonites (cf. Chapter 2) are called “bentones” in the industrial use. This term should not be confused with “bentonite” that refers to the raw clay. In smectites, the layer thickness is around 1 nm, and the lateral dimensions of the layers may vary from 30 nm (Laponite) to several micrometers or larger, depending on the particular clay mineral. The cell parameters describing periodicity in the layers are a ¼ 0.5 nm and b ¼ 0.9 nm. This results in densities of about 2.6 g/cm3. We have mentioned in Section 1.3.1.3 that the basal distance d001 of smectites has an average value of 1.2 nm. In fact, the basal distance corresponds to the sum of layer

FIGURE 1.9 TEM micrograph of a montmorillonite.

4

“Activation” of clays has many meanings in literature; in different contexts, it refers to acid and/or to thermal treatments.

20

CLAYS AND CLAY MINERALS

thickness þ interlayer space and the thickness of the interlayer space is determined by the type of compensator cations located in the interlayer space and their degree of hydration. The value of 1.2 nm is observed for a Naþ smectite under ambient conditions of temperature and water pressure and corresponds to the presence of only one water pseudolayer in the interlayer region; one would therefore expect the d001 value to be dependent on the water activity, that is, on the relative water pressure. Thus, a smectite calcined at high temperature (namely, above 110 C, the temperature chosen as a reference in mechanical studies) will have a d001 value equivalent to those of micas (about 1 nm) since the interlayer will be completely dehydrated; conversely, a fully hydrated smectite shows increasing values until a complete delamination of the clay layers is obtained. At this point, as the long-range order is lost (cf. Section 1.2.2), the d001 XRD peak does not appear any more. The amount of adsorbed water in the interlayer space also depends on the location of the substitution in the layer, as shown by the difference in swelling behavior between beidellite and montmorillonite [27]. The amount of adsorbed interlayer water may be very high and has been underestimated till the 1980s. Moreover, the adsorbed water molecules are present as pillars forming a discontinuous layer more than a continuous layer. When the hydrated cations are ion exchanged with organic cations, in the process of organoclay synthesis, this of course results in larger interlayer spacings (cf. Section 1.4.4.4 and Chapter 2). For clay–polymer nanocomposites (CPN), the d001 peak may disappear as in the case of fully swollen smectites, indicating delamination. 1.3.1.4 Clays, the Oldest Nanomaterials In view of the current fancy for everything “nano,” it is worth underlining that clay structures can rightly be viewed as nanomaterials, since one of the dimensions (namely the thickness of their most basic unit, the layer, along the c axis) is at the nanometer scale. Bidimensional clay minerals can also be considered as inorganic polymers by viewing the repeated half unit cells seen as monomers (Figure 1.10), as already stressed by Bergaya and Lagaly in 2007 [28]. The unmodified clay mineral is a natural nanopolymer of high regularity, which has existed for billions of years on our planet. In fact, the bidimensional phyllosilicates can have one dimension at the nanoscale (for montmorillonite), but also two or three dimensions at the nanoscale as in Laponite or in allophane, respectively. As for the intercalated clays and clay–polymer nanocomposites discussed in the present book, they constitute archetype examples of nanomaterials, but clay nanomaterials have actually been used long before their structure was understood, as witnessed by

FIGURE 1.10 Clays as nanopolymers: a basic unit considered as monomer and repetition of this monomer leading to a structure of the layer.

INTIMATE ORGANIZATION OF THE LAYER

21

the well-known example of “mayan blue” that is a nanocomposite of indigo pigments and clay minerals. This old nanotechnology was successful long before the structure of clay minerals was known. 1.3.1.5 Other Cationic Layered Silicates (T) In the realm of silicates, there also exist alkali silicates (with Naþ as compensating cation) and silicic acids built only from tetrahedral (T) units, for example, the natural minerals magadiite and kenyaite [29,30], whose precise structure is still unknown, and a series of other structures both natural and synthetic [31]. Intercalation of organic compounds into these structures has been largely demonstrated. Studies on their interaction with polymers are few, and most have not been yet evaluated as rubber additives or fillers because of their high synthesis cost; however, magadiite has occasionally been used as a filler [32]. A particular mention must be made of CSH (calcium silicate hydrates) that are an essential constituent of cements. They contain two sheets of Si-containing tetrahedra with a very different organization from those found in clay minerals, surrounding a central sheet containing Ca2þ ions in octahedral coordination [33]. Since the bonding of the Ca2þ to the tetrahedral sheets is rather strong, they could be considered as a “TOT-like” structure. Yet there are important differences in behavior; in particular, CSH do not swell. 1.3.1.6 Nonphyllosilicate Cationic Layered Minerals (O or Mixed T–O) The extended family of layered inorganic materials is not limited to silicates. There exist layers based on negatively charged octahedral sheets (O) with exchangeable hydrated cations between these sheets, for example, layered titanates (Kþ )2(Ti4O9)2  , and so on [34] or titanoniobates (Kþ )(TiNbO5)  , layered manganates, and so on. Zirconium phosphates and phosphonates have mixed sheets containing tetrahedral and octahedral units [35]; they have received considerable attention in the literature since the pioneering work of Alberti and his coworkers [36]. Also layered metal chalcogenides (LMC) involving a broad range of metals, including transition metals, have been the object of many academic and industrial studies due to their interesting intercalation properties [37]. Each of these families may represent a program of intercalation and organophilization in the waiting (initial reports may be found of the use of organophilized titanates as fillers [38]). So far however, the clay minerals family remains the most versatile and the most studied in industrial applications. 1.3.2 Anionic Clay Minerals (O) One particular class of nonphyllosilicate compounds of high interest, is that of anionic clays of the hydrotalcite-like group. Hydrotalcite is probably the only example of a natural LDH. It has been used to synthesize elastomer nanocomposites [39,40]. Its formula unit is Al2Mg6(OH)16(CO3)4H2O. The hydrotalcite denomination probably stems from a perceived similarity with hydrated talc, although the two structures are in fact quite different.

22

CLAYS AND CLAY MINERALS

The layers of hydrotalcite are constituted of a single octahedral sheet with all of the corners of the octahedra being occupied by hydroxide (OH  ) ions. While there exist other minerals such as brucite (Mg3(OH)6) based on uncharged sheets of this type, hydrotalcite and other members of the LDH family contain isomorphous substitutions. Isomorphous substitution in the octahedral sheet leads here to a positively charged sheet (as opposed to the negatively charged layer in cationic clays), and this positive charge is compensated by exchangeable anions in the interlayer. All the other studied anionic clay minerals called hydotalcite-like compounds are synthetic LDH with q 3þ the following general formula: M2þ 1x Mx ðOHÞ3 Xx=q  nH2 O. Their structures involve different types of cations leading to broad nonphyllosilicate families. The interlayer space is very reactive, allowing intercalation of different species, sometimes accompanied by swelling and interesting rheological properties.

1.4 MOST RELEVANT PHYSICOCHEMICAL PROPERTIES OF CLAY MINERAL In the context of the present book, many of the relevant properties of clay minerals are determined by their surfaces and the interfaces they form with other phases. It is important to realize that clay surfaces are chemically very heterogeneous. We will successively consider properties that do not depend on the local chemical features of the surface (cf. Section 1.4.1), and those that do depend on them (cf. Section 1.4.2). 1.4.1 Surface Area and Porosity One of the most common characterization techniques applied to solid materials is low-temperature nitrogen physisorption. Since physisorption is supposedly nonselective, it is generally considered that this technique can provide a quantitative measurement of the surface exposed by the sample, which is then transformed to a specific surface area (SSA or Ss, in m2/g). The SSA is generally extracted from the physisorption data at relatively low pressure using the BET treatment. The physisorption isotherm also gives access to the “porous volume,” and its distribution among micropores (pore diameter <2 nm), mesopores (2 nm
diameter
40 nm), and macropores (diameter 40 nm),5 using standard models such as the Barrett–Joyner–Halenda (BJH) method based on Kelvin’s equation of capillary condensation for mesopores; we refer to standard monographs on the subject for further discussion of the techniques [41,42]. In principle, the experimental value of the SSA should be equal to the geometric area developed by its elementary particles, which can be calculated if the shape and size of the latter is known. Thus, swelling clay minerals and particularly montmorillonite, should have a high total SSA of about 800 m2/g if all the layers are totally 5

These size limits are in agreement with the IUPAC recommended nomenclature.

MOST RELEVANT PHYSICOCHEMICAL PROPERTIES OF CLAY MINERAL

23

exfoliated; on the other hand, we have mentioned in Section 1.2.2 that a SSA value of about 20 m2/g is expected based on the size of the secondary particles in a montmorillonite, and the values effectively measured by N2 physisorption are closer to the second figure. Thus, in general, for TOT clay minerals, two types of surface area (external and internal) must be clearly distinguished [43]. The specific external surface area is the geometric area of the particles (basal and lateral areas of the quasiparallelepipeds constituted by layers stackings) that is accessible to N2 physisorption. The external surface varies from 30 to 130 m2/g depending on the granulometry and aspect ratio of the considered clay mineral, which are related to the number of stacked layers per particle, and therefore to such parameters as the nature of exchangeable cations. The internal surface corresponds to the basal surface of the layers that are “stuck” together within a secondary particle. While it is not normally accessible to nitrogen, it can be revealed by other techniques such as the adsorption of ethylene glycol monoethyl ether (EGME) or glycerol. EGME adsorption [44] was one of the first classical methods used for internal surface area determination; it is based on the idea that EGME forms a bilayer in the interlayer space of clay minerals, as witnessed by the presence of an XRD peak at 1.77 nm. This is equivalent to a single dense layer of EGME molecules on each basal plane, and therefore a measurement of the adsorbed amount will give access to the total basal surface knowing that each molecule occupies a space of 0.44 nm2. However this view of EGME adsorption is somewhat idealized, and in fact its retention depends on the type of compensating cation and layer charge density [45]. In practice, this method only provides a semiquantitative comparison of a series of swelling clay samples [46]. The total surface area (internal þ external) can also be measured by methylene blue (MB) adsorption that, like EGME adsorption, induces swelling. The “MB-spot” is a very rough method used in the field where the amount of swelling clays in the sample is evaluated from the intensity of a blue spot due to methylene blue adsorption. The “MB-titration” in laboratory is more accurate; however, both give similar results [47]. Interesting data on internal and external surface areas can also be garnered using other adsorbing molecules, for example, comparing water and N2 adsorption [27]. The clay surfaces delimit pores with variable sizes (micropores, mesopores, and macropores) and shapes. The hierarchical porosity distribution that can be deduced from the hierarchical arrangement of the clay aggregates/particles/layers described in Section 1.2.1 is not simple to observe. A critical comparison of the numerous techniques of porosity characterization of different porous solids (including pillared clays and LDH) can be found in Ref [48] together with guidelines for the selection of the most appropriate method to follow. For porous clay minerals, lenticular pores are often observed on TEM images. The pores can be accessible to adsorbates or not (open vs. closed pores) depending not only on parameters we have already encountered (chemistry of the clay mineral layer, type of exchangeable cation) but also on the experimental conditions such as air- or freeze-drying [49].

24

CLAYS AND CLAY MINERALS

For macroporosity, the mercury intrusion method remains the main technique of quantitative measurement, even though it has been criticized because mercury is probably able to deform the geometry of the pores during invasion of smaller pores. 1.4.2 Chemical Landscape of the Clay Surfaces The preceding discussion dealt with the geometric extension of the accessible clay surface. The reactivity at a clay interface obviously depends also on its chemical nature and it must be underlined that exposed clay surfaces are chemically heterogeneous. The very existence of a surface may be considered as a defect since it means the interruption of an ideally infinitely extended three-dimensional lattice. An interruption in the c direction results in the exposure of basal planes constituted of siloxane groups (bridging oxygens in Si–O–Si) that are chemically rather inert [50]. On the other hand, termination of the clay lattice in the a and b directions means cutting of partly covalent bonds, for example, Si–O–Si (in the tetrahedral sheets) or Al–O–Al (in the octahedral sheets). This results in “dangling bonds” that in ambient conditions are cured by the formation of, for example, Si–OH or Al–OH2, very similar to those that are found on the surface of nonclay amphoteric oxides such as SiO2 or g-Al2O3. Therefore, the edges of clay particles should exhibit a different reactivity from that of basal planes, one that could be similar to that of amphoteric oxides [51]. As indicated by the term “amphoteric,” a surface group such as Si–OH may react both with a proton (thus giving rise to Si–OH2þ in acidic conditions) or with a hydroxide (thus giving rise to Si–O  in basic conditions). This group may have other chemical properties that are not found on basal planes such as H-bonding or the capacity to act as a ligand; in summary, the edges of clay particles will be able to react in very different ways from the basal planes at solid–liquid or solid–gas interfaces. The amphoteric behavior means that clay minerals have intrinsic acid–base properties; aqueous smectite dispersions have pH values from 7.5 for Mt to 10 for Laponite. The heterogeneity of clay surfaces is also manifested in adsorption from the gas phase. Even for Ar or N2 physisorption that is supposed to be nonspecific, a precise examination of adsorption isotherms at low pressure indicates that the atoms first adsorb on specific sites such as ditrigonal cavities next to Al substitution [52] or more generally sites with high acidity, only later covering the rest of the surface. 1.4.3 Cation (and Anion) Exchange Capacity 1.4.3.1 Definition The notion of cation exchange capacity (CEC) comes from soil science, where it designates the capacity of the soil to hold cations. The soil cations are mostly held by the negatively charged clay minerals (and organic matter particles) through electrostatic forces. As mentioned before, compensating cations must be present in the clay interlayers because the layers bear a negative charge due to isomorphous substitution, and they are easily exchangeable by other cations. For example, a very important mechanism for organoclay preparation is based on the exchange of pristine alkali cations (sodium or calcium) by alkylammonium cations, as will be seen in Chapter 2.

MOST RELEVANT PHYSICOCHEMICAL PROPERTIES OF CLAY MINERAL

25

TABLE 1.3 CEC or AEC Ranges (in meq/100 g) of Some 2:1 Cationic and Anionic Clay Minerals Clay Mineral Species

CEC

Montmorillonite Hectorite Saponite Vermiculite Kaolinite Sepiolite–palygorskite Allophane Imogolite

80–120 120 85 150 3–15 20–30 25 or (10–40) at pH 7 17–40

LDH

AEC 200–400

The CEC represents the total amount of cations available for exchange at a given pH. It is commonly expressed as meq/100 g of calcined clay (meq ¼ “milliequivalent,” i.e., amount in millimole divided by the ionic charge) and can be evaluated by different methods [53]. Typical CEC values of 2:1 phyllosilicates are listed in Table 1.3. Similarly to the CEC of phyllosilicates, LDH present an anionic exchange capacity (AEC). 1.4.3.2 CEC Measurement and Evaluation The CEC is of course directly linked to the charge density of the layers (which itself corresponds to the density of isomorphous substitutions). In first approximation, the CEC is equal to the charge density of the layer (except that it is expressed per unit of mass, rather than per formula unit) if all the compensating cations are susceptible to exchange. The layer charge of smectites and vermiculite groups can be deduced from the chemical structural formula if the sample is very pure and has a known mineralogical composition, or it can be determined by the “alkylammonium method” based on XRD of the organomodified clay mineral, as described by Lagaly [54]. This author has noticed that for smectites the alkylammonium method leads to values 10–20% smaller than the actual CEC. As organic cations are preferentially exchanged over inorganic cations, they have been widely used for semiqualitative CEC determination (and thus for a rough estimate of the amount of swelling clays in pure sample). Particularly, the MB method was used for this purpose as well as for specific surface area determinations. Another method has been applied for dioctahedral smectites with a pronounced beidellitic character, where the tetrahedral substitution that induces the CEC has been estimated through an exchange of the pristine cations by NH4þ cations, followed by the IR observation of these ammonium-saturated clay minerals [55]. As one can guess from the large number of proposed methods, CEC is difficult to determine very accurately, but in any case, its determination requires the complete replacement of all initial exchangeable cations by added “index” cations that should not be present in the

26

CLAYS AND CLAY MINERALS

studied clay. To the best of our knowledge, the most versatile method of CEC determination for several type of clays and clay minerals uses copper complexes with ethylene di- or triamine (e.g., [Cu(en)3]2þ ) as index cations [56]. 1.4.3.3 Cation Exchange Selectivities and Other Chemical Features of Ion Exchange Cation exchange equilibria may be described as simple chemical reactions; however, if one attempts a thermodynamic description, care must be taken to use activities rather than simply concentrations since the participating species are in different phases (solid and solution). Thermodynamic modelization can help to predict the ions exchange, but it should be based on accurate data obtained under fixed experimental conditions. Cation exchange reactions exhibit definite preferences for some cations over others (exchange selectivity). Many data for clay minerals exchange selectivities are available in the literature and they exhibit some regularity. Usually for smectites, there is a preference for larger inorganic cations over the smaller ones, and for cations with higher valence. However, the selectivity of a particular clay mineral for a particular cation is not a simple matter and depends in fact on several physicochemical properties of the cations: their hydration state, their interaction with the clay surface, their polarizability (hard and soft acid base character), and so on, as well as properties of the clay matrix pertaining to their mineralogical structure and including the heterogeneous distribution of charge, the heterogeneity of energy surface sites, site inaccessibility due to aggregation, and so on. As regards the cations, they may be classified in three separate groups: (i) the two monovalent cations Liþ and Naþ , (ii) other monovalent cations with higher radius such as Kþ , and (iii) the bivalent cations. When the exchange reaction on clay minerals is carried out between two cations of different groups, it generally shows a hysteresis effect, which is not yet completely understood. As regards the clay matrix, in some cases, the cations are not easily exchangeable; for instance potassium ions may be trapped in the “ditrigonal” cavities that exist in the tetrahedral layers at the surface of illite and micas, because the radius of Kþ exactly fits the cavity size, in a phenomenon of molecular recognition similar to the selectivity for specific alkali cations of crown ethers. This type of strong selectivity indicates that the macroscopic phenomenon of ion exchange is not entirely explainable in terms of electrostatic interactions between the matrix and the ions; other interaction may be superimposed to it and play a role in determining the energetics of exchange. Although we have approximated the CEC as corresponding to the amount of substitutional layer charges in the above discussion, it is in fact the sum of two kinds of charges: (i) “permanent” charges arising from isomorphous substitutions or site vacancies in the octahedral and/or tetrahedral sheets of the layer, and

MOST RELEVANT PHYSICOCHEMICAL PROPERTIES OF CLAY MINERAL

27

(ii) variable edge charges, induced by the amphoteric behavior of edge groups (cf. Section 1.4.2), which depend on the pH and ion strength of the solution. Their contribution to the total CEC depends of course on the particles morphology and arrangement. The ratio of the edge/permanent surface charge is estimated to lie between 10 and 20% [57] in smectites. In kaolinite particles, on the other hand, there is little to no isomorphic substitution giving rise to permanent charge, and in addition the stacking of layers is more extensive so that the area of edge surfaces exceeds that of planar surfaces. In this case, the CEC is mainly attributed to edge contribution. Finally, it is little known that cationic clay minerals can also adsorb anions in some conditions. The adsorption mechanisms are different from the simple electrostatic compensation of layer substitutional charge; they include the substitution by the anion of variable hydroxyl groups (sometimes called “ligand exchange” because the substituted anion then plays the role of a ligand to the underlying cation). This occurs at the edges of TOT clay minerals, and also on basal surfaces in the case of kaolinites. It is possible to modify the CEC values of a clay mineral by applying what is called the Hofmann–Klemen (HK) effect [58]. Clay minerals are first partially exchanged with small cations (generally Liþ ) and heated at about 200–300 C, followed by glycerol adsorption. The thermal treatment leads to an irreversible migration of the small cations inside the layer. This apparently irreversible ion fixation decreases the layer charge and thus the CEC. While the HK effect has a respectable history, its mechanism was recently revisited and there is some controversy concerning the location of the migrated cation after heating and the reversibility of this effect. 1.4.4 Intercalation and Confinement in the Interlayer Space 1.4.4.1 General Definition Most of the technological applications of clays are related to reactions that occur in the interlayer spaces. The properties of “intercalated” species (i.e., species present in the clay mineral interlayer) are often different from those of the same species in solution; the interlayer may be considered as a nanoreactor that imposes on the intercalated species several modifications collectively known as confinement effects, which may be due to local chemical environment, size restrictions, specific orientation, mobility limitations, and so on. Confinement is not limited to clay minerals, it may also be observed in zeolites, mesoporous silicas, in the cavities of protein enzymes, and so on, but clay minerals are specific in providing tunable cavity size because of the variable interlayer distance. 1.4.4.2 Intercalation of Small Polar Molecules (e.g., Water) Intercalated water, naturally present in clays, has a very high acidity compared to aqueous medium. In the interlayer space, polar molecules such as water or alcohols interact with compensating cations through cation–dipole interactions, which no doubt have some amount of covalent character and could therefore be called coordinative bonds; this results in a weakening of the O–H bond and thus increased acidity.

28

CLAYS AND CLAY MINERALS

This type of intercalation mechanism may also be called “inner-sphere adsorption” because the water molecules are in the coordination (i.e., inner) sphere of the cations. Further water molecules are not directly bound to the cation but form H-bonds with the first sphere, in “outer-sphere adsorption.” By using the desorption of a small polar molecule, namely methanol that contains a single hydroxyl group as opposed to the two O–H of water, it was shown [59] that in addition to the welldefined coordinative interaction with different compensating cations, the adsorbed molecules also form hydrogen bonds between themselves in the interlayer space and perhaps even to the siloxanes of the basal clay surface. Similar phenomena may occur with other polar solvents. In some cases, a single clay particle may contain solvent-filled interlayers of different thicknesses. This leads to interstratified phases [60–63] that have been mentioned before: alternance of zero, one, two, or three water pseudolayers (or water pseudosheets) between the TOT layers (strata). 1.4.4.3 Intercalation of Bulky Inorganic Oligomers Some bulky cationic species such as the [Al13O4(OH)24(H2O)12]7þ ions with anti-Keggin structure, called “Al13” ions for short, may lead to intercalated clay minerals by ion exchange of the initial compensating cations. After heating, these oligomers become irreversibly fixed between two adjacent layers and act as “pillars” leading to microporous clay minerals with constantly accessible interlayer porosity called pillared clays or PILC [64,65]. This permanent pore opening could favor the intercalation of other species in the empty microporosity. Although these materials have been extensively studied in literature, and despite many hypotheses proposed to explain the fixation of these pillars in the interlayer space after heating, the pillaring process is not well understood and there is probably no universal mechanism for this irreversible polycation fixation. 1.4.4.4 Intercalation of Organic Species There are several routes to the intercalation of organics corresponding to different intercalation mechanisms. Neutral organic ligands can form complexes with the interlayer cations, and more generally polar organic solvents can displace water. Alternatively, the pristine cations can be exchanged by organic cationic species. The most studied, and most relevant example is the exchange by alkylammonium cations (most often quaternary ammonium). This will be developed in more detail in Chapter 2 since it is the starting point of the modified bentonites called bentones. Intercalation by grafting reactions is also often seen in literature. Here covalent bonds established by condensation of some reactive groups on the organic chain with the clay silanol surface groups or with the hydroxyl groups of octahedral layer of kaolinite allows the expansion of the interlayer space. Consequently, the surface becomes hydrophobic. This route will also be treated in more detail in Chapter 2. 1.4.4.5 The Particular Case of Intercalation into Kaolin It is quite frequent to use preswelled kaolin in CPN application, since this clay is the

MOST RELEVANT PHYSICOCHEMICAL PROPERTIES OF CLAY MINERAL

29

least expensive on the market. Three main categories of molecules can give intercalates in kaolin: . . .

Compounds like hydrazine and formamide that are both H-bond acceptors and donors and can swell the interlayers by forming H-bond networks. Compounds like dimethyl sulfoxide (DMSO) with high dipole moments. Compounds like salts of short-chain fatty acids.

1.4.4.6 Intercalation of Polymers Of particular concern for the present book is polymer intercalation in swelling clay minerals. Many polymers are able to directly intercalate with pristine clay minerals, but they are mostly restricted to hydrophilic polymers bearing, for example, ether or alcohol groups: PEO (polyethylene oxides) [66,67], PVA (polyvinyl alcohol) [68], PEDOT (poly(3,4ethylenedioxythiophene)) [69], polypeptides or proteins such as gelatin [70], polylysine, and polyglutamic acid [71], and so on. For a recent update, see Ref. [72] (especially Section 3.9 in this reference). Hydrophilic polymers/clay nanocomposites can be used instead of more classical organoclays as rubber fillers. In practice, in many reported syntheses, polymers are probably adsorbed on the external surface of the particles rather than intercalated. However, if the starting clay mineral is delaminated and well dispersed, nonionic polymers can definitely be intercalated from aqueous solutions. The basal distance observed is usually not very large and corresponds to a thickness of one or two linear macromolecule chains at maximum. The picture of clay–polymer interaction is certainly complex, showing simultaneous loops of polymer chain that remain amorphous in the interaggregates porosity and part of polymer tails intercalated in the interlayer space [71] and leading to more crystalline domains due to the nanoconfinement of the intercalated part of polymer. The existence of cationic groups on the polymer, such as in polylysine, is a particularly favorable factor for intercalation into smectites since the electrostatic interaction with negatively charged layers will provide a driving force for intercalation. For neutral polymers and/or neutral layers (e.g., in kaolinite), the formation of intercalated layer does not occur by ion exchange but by other interactions such as H bonds, ion–dipole interaction, charge transfer, and van der Waals forces. All of these interactions contribute to the enthalpic part of the driving force for polymer intercalation. As regards the entropic part, it is more difficult to evaluate; the intercalation of a single polymer molecule is expected to displace many interlayer solvent molecules, which should be favorable, but the effects on the configuration entropy of the polymer are not so clear. Polymers intercalation into pristine clay minerals is usually carried out from solutions, although there are a few successful examples of intercalation from polymer melts [73,74]. This procedure will be discussed later in the present book in the frame of organoclay–polymer nanocomposites; in spite of its industrial importance, it has probably not received enough attention in academic studies so far. It should be noted, however, that even authors reporting direct polymer intercalation in unmodified clays state that apolar hydrocarbon polymers cannot be intercalated in this way.

30

CLAYS AND CLAY MINERALS

A completely different procedure to obtain clay/intercalated polymer nanocomposites is the in situ polymerization of previously intercalated monomers such as acrylonitrile [75] leading to intercalated polyacrylonitrile (PAN). This often occurs through enhancement by an initiator molecule or UV radiation or temperature. In summary, intercalation of inorganic and organic species in clay minerals is a vast research domain. Intercalation can occur from vapor, from solutions, from pure melt, and in the solid state. The latter two procedures that are the most interesting for industrial use (avoiding the use of polluting solvent) are very scarce in academic studies. 1.4.5 Swelling The ability to swell in water (or in any polar solvent) is a defining property of the smectite group of cationic clay minerals. Swelling is easily observable at the macroscopic level, by the formation of gel-like phases when water is added to clay minerals, with a large increase in volume as compared to the dry solid. A swelling index may be defined as the gel volume (in mL) obtained per 10 g of clay mineral, and it may reach values of 25–30 mL/10 g [76]. There is some ambiguity here because “swelling” is used to denote a macroscopic phenomenon as well as its molecular counterpart. Swelling as noticed at the macroscopic level, with gel formation, betrays some kind of microscopic scale structuration of the solvent–clay system. In principle, this could simply be a consequence of the formation of associations between the clay particles and aggregates (cf. Section 1.2.1). However, when used at the molecular scale, this term most often refers to the intercalation of large amounts of water in the interlayer space between the layers resulting in increasing interlayer distances. Swelling is therefore a process whose end result may be high delamination of the clay layers and even exfoliation, if the distance between the layers is large enough. The first steps of swelling are discrete, as they consist in the intercalation of crystallographically ordered water layers; discontinuous jumps in the d001 spacing are then observed with increasing water content [77]. When additional water is intercalated, at a given point, intercalated water molecules are not rigorously ordered any more, and therefore swelling is better described as a continuous phenomenon; osmotic swelling, that is, migration of water from the bulk solution into the interlayer region, driven by the lower water activity there. This osmotic transition (OT) is manifested by the transition from a hydrated powder to a paste. If more water is added, the clay dispersion remains in the form of a gel that can persist down to low overall clay contents: about 10% for Mt in water or as low as 3% for Laponite in water (a beaker containing a 3% Laponite dispersion can be turned upside down without spilling its content). Upon further dilution, the sol–gel transition occurs at the “SGT” threshold, where isolated aggregates are formed. In even more dilute dispersions, isolated particles are formed after a “local osmotic transition” (LOT) that is found at about 3% for Mt in water [78]. Whether swelling proceeds to delamination or not hinges on a delicate balance between the osmotic forces and the electrostatic interactions, which depend on the

MOST RELEVANT PHYSICOCHEMICAL PROPERTIES OF CLAY MINERAL

31

compensating cations charges, and the layers charge density; for instance, Naþ montmorillonites can swell extensively. It would seem that at an intermediate degree of swelling, the morphology of layers stackings goes from “columnar” to “zigzag,” for example, in the case of kaolinite, where the kaolinite layers become free to slide alongside each other. Recent research reveals additional phenomena, thus, in the gel region, smectite dispersions can organize as nematic liquid crystals and exhibit a nematic-to-isotropic transition occurring before the sol–gel transition [79,80]. The swelling behavior of clay minerals is important to consider in the frame of CPN because clay interaction with surfactants or with water-soluble polymer occurs in aqueous media. 1.4.6 Rheology Smectites display a complex rheological behavior when dispersed in an aqueous medium. When present in low concentrations they form dispersions with almost Newtonian properties. In other words, the consistency curve of such dispersions _ in s  1) is linear, the slope of the straight line (shear stress t, in Pa, vs. shear rate g, giving access to the dispersion viscosity. At higher clay concentrations, a Bingham plastic behavior is observed, that is, the consistency curve becomes linear only after a threshold shear stress is exceeded, the so-called yield stress, t e [81]. The rheology of clay dispersions is a testing ground for competing models of interaction in colloid science, and we will only offer some basic developments here. The viscosity of inorganic dispersions is classically explained by the fact that the particles hinder each other’s motions, and distort flow patterns in their vicinity, leading to increased friction. Einstein proposed that the viscosity of a dispersion containing a size fraction f of dispersed particles is given by ðh=h Þ ¼ 1 þ ð5f=2Þ [82,83], where h is the viscosity of the dispersion and h is the viscosity of the pure liquid. Actually, in a dispersion of charged particles, the effective size is not that of the particle itself, but of the particle surrounded by its “diffuse layer,” and electroviscous effects are at work. This needs some explanation. Since electric charge cannot accumulate macroscopically, a charged particle in an electrolyte solution will induce the formation in its vicinity of a swarm of ions of the opposite charge; for instance, the basal planes of a smectite clay particle that bear a net negative charge will be surrounded by a region of the solution where the positive ions (Naþ , Ca2þ , etc., as the case may be) outnumber the negative ions. The thickness of this diffuse swarm or diffuse layer extends to a few nanometers (high ion strengths cause diffuse layer contraction), and the properties of the “double layer” are extensively studied in colloid chemistry (e.g., Chapters 3 and 5 of Ref. [84]), as well as their effects of interparticle interactions (the positive diffuse layer of any clay particle will repel that of its neighbors). The “DLVO” theory (Derjaguin–Landau–Verwey–Overbeek, see, e.g., Chapter 8 of Ref. [84]) explains the stability of colloidal dispersions by the interplay between double-layer repulsions and van der Waals attraction.

32

CLAYS AND CLAY MINERALS

When the dispersion is submitted to shear flow, the diffuse layers will tend to oppose the ensuing deformation, causing further friction upon the liquid, this is called the “first electroviscous effect,” while the “second electroviscous effect” refers to the viscosity increase due to the electrostatic repulsion between colloidal particles during flow-induced encounters. These considerations can explain the variations of viscosity in the Newtonian regime. As regards the existence of a threshold stress (yield stress), it is correlated with a long-range organization of the dispersion that we have already mentioned when discussing the stability of clay gels (Section 1.4.5). The effect on yield stress of varying liquid phase parameters such as pH and ion strength has been studied and it is not straightforward; for example, at high clay concentrations an increase in the ion strength causes an initial decrease of the yield stress, followed by a slow increase. As another example, for a Naþ -montmorillonite, there is a maximum in shear stress (at constant shear rate) at pH 7, but it disappears when the compensating cation is Ca2þ ; and the behavior of the dispersion can change from Newtonian to Bingham plastic as the pH is varied ([85], pp. 204–205). This is because the structure of the dispersion will be the result of different types of interparticle association: face-to-face, face-to-edge, edge-to-edge (cf. Section 1.2), each one reacting independently to changes in those parameters, and each one able to give rise to clay particles aggregation. The chemical heterogeneity (or patchwise heterogeneity) of the clay surface underlined in Section 1.4.2 is particularly important here. Recall that basal planes of smectites bear a permanent, pHindependent negative charge; in contrast, the edge groups of clay minerals (and the exposed OH-bearing octahedral sheets of kaolin) can be positively or negatively ionized by pH-dependent reactions. As a result, the latter regions of the clay surface will have a net positive charge at low pH (where, e.g., Si-OH2þ groups predominate), and a net negative charge at high pH (where Si-O groups predominate). Thus, in some conditions the edges can have a charge of opposite sign to the basal faces, and it is not hard to understand that this will influence the preferred type of particle association. Further interesting developments may be found in the work of Tombacz and Szekeres for montmorillonite [86] and for kaolinite [87]. For these amphoteric surfaces, there must be a pH value at which positively charged groups will be present in exactly the same number as negatively charged ones. This is called the point of zero (net protonic) charge or PZC (PZNPC). The PZC is important to understand the colloidal properties of a dispersion and can be measured simply by protonic titration; for example, a PZC of 6.5 was measured for montmorillonite edge surfaces [86]. The PZC must not be confused with the isoelectric point (IEP) obtained from electrophoretic mobility measurement. The IEP is the pH for which a particle will not migrate in an electric field, and it is determined by the overall charge of the particle (basal planes þ edges, which may be of opposite sign), and in addition the particle usually moves in a solidary way with a few water layers that will contain part of the compensating cations in the diffuse layer. It may thus happen that a well-defined PZC may be measured, although the system as a whole does not have an IEP. More generally, it has been proposed to relate the zeta potential (calculated from electrophoretic mobility data measured as a function

AVAILABILITY OF NATURAL CLAYS AND SYNTHETIC CLAY MINERALS

33

of pH) to clay dispersibility. Although qualitative, and in some cases, semiquantitative agreement between the zeta potential of clay dispersions and their dispersibility has been reported, the usefulness of zeta potential measurements as an indicator of the dispersibility of colloidal systems has been strongly contested [88,89]. Finally, rheological properties may change upon mixing, and also simply upon aging, often in a reversible manner (thixotropy) [90]. There is continued interest in their evaluation due to the relevance for such phenomena as mud flows, or applications to drilling fluids. The relationship between bentonite dispersion/flocculation and the rheological behavior of the corresponding dispersions, their dependency on aqueous phase parameters, and the various theoretical attempts to explain this dependency have been the object of a review by Luckham and Rossi [91].

1.5 AVAILABILITY OF NATURAL CLAYS AND SYNTHETIC CLAY MINERALS Natural clays are very abundant, even though clays and clay minerals form under a fairly limited range of thermodynamic conditions. The environments of formation include soil horizons, continental and marine sediments, geothermal fields, volcanic deposits, and weathering rock formations. Most clay minerals form where rocks are in contact with water, air, or steam. Examples of these situations include weathering boulders on a hillside, sediments on sea or lake bottoms, deeply buried sediments containing pore water, and rocks in contact with water heated by magma (molten rock). All of these environments may cause the formation of clay minerals from preexisting minerals by dissolution– reprecipitation. The soil solution concentration of Si and Al from the weathering of volcanic rock is one of the most important factors determining the type of clay minerals that will be formed. For instance, the Al:Si ratio of allophane ranges from 1 to 2 according to the conditions of formation. They have been identified as being of three types: (1) Al-rich allophane (proto-imogolite), (2) Si-rich allophane, and (3) an hydrous allophane (feldspathoid allophane) [92]. Extensive alteration of rocks to clay minerals can produce relatively pure clay deposits that are of economic interest (e.g., bentonites used for drilling muds and clays used in ceramics). In addition to the weathering process, some clay minerals are formed by hydrothermal activity. Clay deposits may be formed at places as residual deposits in soil, but thick deposits usually are formed as the result of a secondary sedimentary deposition processes after they have been eroded and transported from their original location of formation. Primary clays, mostly kaolins, are located at the site of formation. The already mentioned bentonites form mainly from alteration of pyroclastic and/ or volcanoclastic rocks. Extensive deposits, linked to large eruptions, have formed repeatedly in the past. Bentonites generally form by diagenetic or hydrothermal alteration, favored by fluids that leach alkali elements and by high Mg content. Smectite composition is partly controlled by parent rock chemistry. Recent studies

34

SiO2, FeC12, Mg-, Al-salts

SiO2, Al(OH)3, Mg(OH)2, NH4OH (or NH4F) Na2SiO3, metal salts, acid

Gel from Na2SiO3, NaAlO2 in HCl Gel SiO2, Fe(SO4)2

TEOS, MgCl2, CuCl2, NaOH, LiF, NaF

TEOS, AlCl3, FeCl2, Ca(OH)2, hydrazine

Mg(OH)2,organic cations, LiF

TEOS, Mg(NO3)2, Al(NO3)3, Na2CO3, Li2CO3 SiO2/Al2O3 gel, divalent metal nitrates, urea Silica sol (Ludox), Mg(OH)2, TEA, LiF

TEOS, AlCl3, FeCl2, Ca(OH)2, hydrazine

SiO2–Al2O3–MgO (ZnO) –Na2O H2O SiO2, Al2O3, MgO, HF, H2O Silicic acid, MgCl2, Na-, K-, Ca-hydroxide

SiO2, FeCl2, Mg-, Al-salts SiO2, Al(OH)3, Mg(OH)2, NH4OH (or NH4F)

[95–99]

[100,101] [102]

[103,104] [105,106]

[107]

[105,108,109]

[110,111]

[112,113]

[116–121]

[122] [123] [124]

[125] [126]

[114] [115]

Silicic acid, MgCl2, Na-, K-, Ca-hydroxide

Starting Materials

[94]

References

Boiling 100 C, in very dilute solutions 300–350 C

95 C, 15 days, redox cycles, 10–13 weeks 220 C, 72 h 220 C, 48 h Boiling, 20 h

90 C, <2 days, homogeneous solution 100 C,
3days

400 C, 1 month

100 C, reflux, 2 days

<100 C,
8 weeks

Boiling 45 days in Mg acetate 75 C, 15 day, 1 month; 100 C, 1 month; 150 C, 12 days Reflux, 12 h

Boiling at 100 C, in very dilute solutions 300–350 C <100 C

Boiling, 20 h

Conditions

Obtained Phases

Badly crystallized saponite Hectorite, crystallinity and stacking increase with TEA Hydroxy-interlayered, Fe-beidellite, poorly crystalline Montmorillonite Saponite Amorphous hectorite, small amount of poorly crystalline Ht Hectorite, Fe-saponite, nontronite Saponite

Saponite Hectorite, saponite, stevensite, phases depend on starting composition Saponite Nontronite, trace-Fe(OH)2; Fe-smectite under oxidizing conditions (“defect” nontronite) Cu-, Mg-fluorohectorite possible total replacement of Mg by Cu Saponite, nontronite formed phases depend on starting composition Organo-hectorite, direct uniform intercalation, but crystallinity varied with ion Saponite

Amorphous hectorite, small amount of poorly crystalline Ht Hectorite, Fe-saponite, nontronite

TABLE 1.4 Some Examples of Procedures for the Obtention of Synthetic Clays in Moderate Conditions

CLAYS AND (MODIFIED) CLAY MINERALS AS FILLERS

35

have shown that bentonite deposits may display cryptic variations in layer charge— that is, the variations are not visible at the macroscopic scale—and these correlate with physical properties. As an unfortunate consequence, the properties of clay minerals obtained from the same deposit may show some variability as the resource is used up and different parts of the deposit are dug and collected. Geographically, clays are not a scarce resource as clay deposits are spread all over the word (Algeria, Morocco, Tunisia, Spain, Sardinia, Greece, Russia, China, USA, etc.). They differ by their composition and the nature of the impurities. They are available in big quantities and most deposits are commercialized. Two major problems encountered in clays exploitation and industrial use are the depletion of the deposits, at least those with easy exploitation access, and their occurrence as a mixture of several phases, instead of a pure clay mineral phase. To overcome these inconvenients, laboratory syntheses of clay minerals have been developed [93]. A lot of effort has been deployed to obtain well-crystallized clay minerals with controlled composition and properties, starting from initial oxides. Usually, synthetic clay minerals are crystallized from gels under hydrothermal conditions. Several methods have been reported in the literature where a great number of parameters have been changed such as the nature of the starting gel, the temperature, the pressure, and the duration. Some relevant results are reported in Table 1.4. All these syntheses were performed at temperatures ranging from 150 to 450 C and pressures ranging from the autogenous water pressure to 1500 bar. The syntheses were performed in alkaline or fluorine medium. Trioctahedral clay minerals such as hectorite and saponite that are based on Mg oxides as starting materials are the easiest to synthesize. Laponite, first commercialized by Laporte and more recently by Tolsa,6 is a synthetic trioctahedral hectorite with very small size (its particles are discs with a diameter of about 30 nm). Dioctahedral species as montmorillonite are more difficult to obtain in laboratory and are usually very ill-crystallized. Until now, the conditions required for clay minerals synthesis make their commercialization very expensive and thus limited to high-added value applications. Therefore, so far, most commercial applications are based on natural clays after the application of several purification steps to eliminate the major impurities, even though synthetic clays can technically be used as filler materials for CPN [127].

1.6 CLAYS AND (MODIFIED) CLAY MINERALS AS FILLERS Clay minerals are a key material in a great number of agricultural and industrial applications [128]. Particularly, they are used as fillers in the formulation of various products based on polymers matrices as, for example, in drilling fluids, in the paper industry, and/or in rubber. One of their advantages is that clays are attractive in our time of ecology-conscious chemistry; most probably, unforeseen adverse health effects (such as became painfully obvious for asbestos) need not be feared in the case of clay minerals that have accompanied mankind for millenia—for instance, natural 6

Laponite (Laporte, Rockwood) and Fluorohectorite (Somasif ME 100).

36

CLAYS AND CLAY MINERALS

montmorillonite is used as an oral drug, sold in pharmacy under the trade name of “Smecta.” The first reason for industrialists to add nanoclay fillers to polymers was simply to reduce costs as compared to previously used fillers such as calcium carbonates or glass fibers. But it later turned out (actually it was not realized until the last two decades) that a small amount of clay added as a filler could bring additional benefits, as exemplified the synthesis of nylon 6–clay hybrid [129] that has been commercialized by Toyota in Japan. In fact, in other domains it was already known that their small size (typically <1 mm, see Section 1.1) and large aspect ratio allow clay minerals to strongly influence the physical properties of soils and sediments. The paramount importance of clays in determining the CEC of soils has been mentioned before; they are also used as pond liners and for the isolation of hazardous wastes. Just as in soil science and in geomechanics, small amounts of clay mineral can dramatically influence the final properties of CPN in nanotechnology. A few percent of clay minerals are incorporated in various CPN to improve the initial properties of the major (polymer) component or to gain new properties not seen in the net polymer alone. Several important books, reviews, and book chapters have been dedicated to these CPN [130–135]. The superiority of clays over other nanofillers has been reported many times in the literature and particularly the high performance of elastomers using clays as nanofillers. Among the vast family of layered silicates, Mt is particularly attractive as reinforcement for CPN because it is environmentally friendly, readily available in large quantities with a relatively low cost. Moreover, its intercalation chemistry is more or less well understood. Naþ -Mt is hydrophilic and swells readily when immersed in water. Via cation exchange of Naþ with alkylammonium ions under proper conditions, the clay mineral surface can be converted from hydrophilic to organophilic. The space between the silicate layers depends greatly on the length of the alkyl chain and the ratio of cross-sectional area to available area per cation. The conversion of hydrophilic inorganic clay to a hydrophobic organoclay also improves the interfacial adhesion properties between the organic and inorganic phases when a hydrophobic polymer matrix is involved. Nowadays, several pristine clays and organoclays are available commercially at relatively low cost, and are used in different CPN applications. In conclusion, clays, clay minerals, and organoclays are actually attracting exponential research interest and seem to be promising CPN fillers for applications in the future. Their properties, although intensively studied, are not yet completely understood, particularly concerning the interaction with guest species (water, surfactants, polymer, etc.) at the clay interfaces. Fundamental clay research in all disciplines still leads to open questions. For example, and particularly relevant for the subject of this book, it is difficult to rationalize and even more to foresee the interaction mechanisms of a given clay with a given polymer, and the behavior of the partially intercalated polymer in the confined interlayer space is not very clear. Nevertheless, clays are an abundant resource with a high potential for the preparation of a wide range of high-performance nanocomposites.

REFERENCES

37

The huge potential applications of organoclays as effective reinforcement for polymers have prompted many researchers to investigate their structural behavior. The next chapter will concern the different modifications of clay minerals by organotreatments and other procedures in order to use them as fillers in polymers to obtain CPN.

ACKNOWLEDGMENT The authors are particularly grateful to Dr. Fabrice Gaslain for preparing the structural figures.

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15. 16.

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CHAPTER 3

INDUSTRIAL TREATMENTS AND MODIFICATION OF CLAY MINERALS CINZIA DELLA PORTA

3.1 BENTONITE: FROM MINE TO PLANT 3.1.1 A Largely Diffused Clay Section I of this book gives a description of layered silicates, both synthetic and naturally occurring and, in the first two chapters, clays and organically modified clays are in detail discussed. This chapter covers treatments and modifications, in an industrial environment, of an inorganic clay, from mining to purification and modification, to reactions with an organic modifier. As most of the rubber–clay nanocomposites (RCN) are based on cationic clays, this chapter focuses its attention on a worldwide largely diffused cationic clay: bentonite. Bentonite is a clay, from the family of smectites, with a significant content of montmorillonite: these clays are the most used for the preparation of rubber–clay nanocomposites. As discussed in more detail in Section 3.1.2, the origin of bentonite is the alteration of volcanic ash or the hydrothermal alteration of volcanic rocks. The structure of bentonite consists of two types of layers: silica tetrahedral layers and aluminum octahedral layers, the two tetrahedral layers sandwiching the octahedral one. These three layers form one clay sheet, where an isomorphous substitution of aluminum(III) by magnesium(II) can occur. On the clay surface, the negative charge is compensated by cations, whereas the positively charged clay crystal edges are compensated by anions. The elementary particle is a platelet, about 1 nm thick, with lateral dimension extending to about 1 mm. The various platelets are held together by electrostatic and van der Waals forces, but they can be dispersed in submicronic particles until a specific surface area of 800 m2/g is developed. This capacity, to show a so large specific surface area, is one of the cause of the great water absorption of this material. Waterproofing properties of bentonite are indeed given by water absorption as well as by the clay swelling. Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.  2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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Classification of a bentonite is made by first investigating structural characteristics such as its chemical and mineralogical composition, the cation exchange capacity and the specific surface area, whereas colloidal and binding properties have to be assessed in order to define the possible clay applications. It is in fact known that naturally occurring bentonites are either sodium– or calcium–bentonite and that the class of sodium–bentonites is made of clays endowed with different properties. Worldwide bentonite consumption has been steadily increasing since early 1900s, as bentonite takes advantage of the structural differences among its various types to allow many different and spread applications. Na–bentonite is characterized by its ability to adsorb large amount of water and form viscous, thixotropic suspensions. Ca–bentonite, in contrast, is characterized by its low water absorption and poor swelling capability and by its inability to stay suspended in water. Each type of bentonite has thus its own unique application. Table 3.1 summarizes commercial applications of a bentonite as a function of its physicochemical properties.

TABLE 3.1 Commercial Applications of Bentonite in Relation to Its Physicochemical Properties Physicochemical Properties Sorptive properties

Surface area

Rheological properties

Impermeability, coating properties Bonding properties

Plasticity

Commercial Application—What Bentonite Does Refining and bleaching of oil Clarification and purification of wine, juices, sugar solution Water purification and effluent treatment Paper production Pharmaceutical and therapeutic preparation Adsorbent pet and animal litter Catalytic action, carrier for catalysts Carrier for insecticides Mineral filler and extender Nanoclay Drilling fluids Paints (oil and water based) Fertilizer sprays Bitumen emulsions Formulation for ceramic glazes Wall support for boreholes Civil engineering (diaphragm wall construction) Nondrip paints Impermeable membrane Drilling in permeable strata Bonding foundry molding sands Pelletizing iron ore concentrates Pelletizing animal feedstuffs Formulation of mortars, putties, adhesives, some ceramic bodies

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89

In the light of what reported so far, an important step of the production process for the preparation of a modified clay is the selection of an ad hoc type of bentonite: a naturally occurring one or, for example, a sodium activated one. 3.1.2 Geological Occurrence Most naturally occurring bentonites were formed by alteration of igneous material, a process that led to two markedly different types of bentonite deposits: one resulting from subaqueous alteration of fine-grained volcanic ash and another resulting from in situ hydrothermal alteration of acid volcanic rocks [1,2]. The alteration of volcanic pyroclastic rocks to smectite is a common process, often producing bentonite deposits of higher commercial value. Deposits of economic grades of bentonite may be formed by alteration of igneous material by of the following mechanisms: (1) diagenetic alteration of fine-grained volcanic ash and (2) hydrothermal alteration of volcanic rocks [3]. Regardless of the formation mechanism, leaching of alkali elements and a high Mg2 þ /H þ ratio are required to form smectites instead of zeolites during the alteration of volcanic glass [4]. Bentonites range in color from black through to white but most frequently are bluish-green when fresh, weathering to a yellowish-brown color at outcrop,1 due to the oxidation of ferrous iron. Clays collected in the neighborhood of an outcrop often exhibit enhanced swelling properties. Despite their usual characteristic appearance at outcrop, where they tend to exhibit a frothy or popcorn texture due to successive wetting and drying, deposits may be easily overlooked during field mapping. In addition, because bentonite deposits sometimes exhibit a highly lenticular form, some may have no surface expression with no outcrops. Detection then can prove extremely difficult. 3.1.3 Mining Bentonite is always extracted in open pit working,2 usually by opencast methods in which the overburden3 is dumped into the worked out sections of the pit and rapidly restored. The relative proportion of overburden to mineral is the overburden ratio, it may be 20 : 1 for high-quality bentonite or 5 : 1. Variations in the properties of bentonite can evidently occur within the same bed,4 as well as in different beds and this means that selective mining is often necessary to match different qualities with specific applications [5–7]. Bentonite is usually extracted by surface mining as shown in Figure 3.1. Bulldozers and tractor-scrapers are used to remove overburden material, topsoil

1

An outcrop is a visible exposure of bedrock or ancient superficial deposits on the surface of the Earth. Open-pit working (or mining) refers to a method for extracting rocks or minerals form the earth by their removal from an open pit, that is, from a hole in a surface. 3 An Overburden is the material that lies above an area of rocks of economic or scientific interest. 4 Bed: is the bentonite layer in the mine. 2

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INDUSTRIAL TREATMENTS AND MODIFICATION OF CLAY MINERALS

FIGURE 3.1 Bentonite mine.

and subsoil5 are stockpiled separately. After the clay has been extracted from the pit, the overburden, subsoil and topsoil6 are replaced.

3.2 PROCESSING OF BENTONITE 3.2.1 Modification of Bentonite Properties Major bentonite product groups are Ca–bentonite, Na–bentonite, Na-exchanged bentonite, acid-activated clays, organoclays. Following removal from the mine, bentonite is hauled to a processing plant. Processing technologies depend both on the nature of the crude bentonite and on its final application [8,9]. In fact, bentonite is rarely used in the raw form and undergoes processing essentially with the aim to modify its properties in view of specific industrial applications. In general, processing is designed with the aim to obtain a more efficient clay dispersibility, through (i) increase of clay surface area, (ii) modification of surface chemical properties, (iii) increase of montmorillonite content. As far as the surface area of bentonite is concerned, the highest values are achieved when montmorillonite is fully exfoliated to give individual sheets. Mechanical shear is one method used to exfoliate montmorillonite. This is typically accomplished either by high-speed mixing, when the bentonite is in 5 6

Subsoil is the layer of soil under the topsoil on the surface of the ground. Topsoil is the upper, outermost layer of soil, usually having a thickness ranging from 5 to 20 cm.

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91

fluid state, or by high-pressure extrusion, when bentonite is in a semisolid state. Exfoliation may also be facilitated by chemical modification of montmorillonite surface through cation exchange. The water stability and colloidal stability of a clay are improved when exchangeable divalent cations, Ca2 þ and Mg2 þ , are replaced by univalent cations, Na þ and Li þ . In one common approach, Na2CO3 is added to bentonite to produce a Na-rich clay. Sodium exchange improves dispersibility in water, whereas it is of little benefit in weakly polar and nonpolar liquids, such as toluene or mineral oil: in this case, the driving force for clay dispersion is largely entropic rather than enthalpic. Where dispersion of clay in nonpolar fluid is required, hydrophilic Na–bentonite are transformed to hydrophobic organobentonites by exchanging the surface cations with positively charged organic species [10]. 3.2.2 Processing Technologies Common bentonite processing steps can to include extrusion, drying, milling, screening, air classification, centrifugation, agglomeration, acid leaching, and cation exchange. In Scheme 3.1, there is a simplified flow diagram for the processing of a Ca–bentonite. Initial processing steps consist of drying and grinding.

Raw bentonite from the mine

Treatment with Na2CO3– activation step

Sieving

Drying

Milling—classification

SCHEME 3.1 Block diagram process of a raw Ca–bentonite.

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INDUSTRIAL TREATMENTS AND MODIFICATION OF CLAY MINERALS

The drying step is designed in the light of two main requirements: (i) To preserve the structure of the clay. (ii) To maintain the clay moisture level at about 10–13 wt.%, that is, at the equilibrium level. In particular, it is of the outmost importance not to damage any part of the clay structure, to allow the clay itself to develop its typical colloidal and binding properties. To this purpose, clay cannot be dried at too high temperature, that has to be maintained, in the oven, under 300 C. The experimental conditions of the drying step are tuned as a function of the smectite clay structure, that is investigated through a differential thermal analysis (DTA). DTA allows in fact to characterize the smectite constituent of a bentonite and to identify impurities. DTA curves of smectite are usually divided into three regions: (i) Low temperature region (<300 C) within which adsorbed and exchangeablecation coordinated water are released. (ii) Dehydroxylation region (400–750 C). (iii) High-temperature region ( > 800 C), where new phases crystallize from the dehydroxylated clay. In the low-temperature region, smectite containing monovalent exchangeable cations show only a single endothermic peak, whereas those containing divalent Bentonite

Water dispersion

Purification

Reaction

Drying/milling

SCHEME 3.2 Block diagram nanoclay production.

PURIFICATION OF CLAY

93

exchangeable cations show a well-defined shoulder to this peak at 200 C representing expulsion of water loosely bound to these cations (the “hydration shell”). DTA findings give thus information on the clay composition and confirm, as anticipated above, that the drying step temperature has to be maintained below 300 C, to avoid any transformation of the smectite structure. The dried clay is then either screened and marketed in the granular form or milled to fine powder, size grading being carried out by air classification. Fines powder may be pelletized and crushed to produce a higher yield of granules. Granules may be lightly calcined to make them water stable. In Scheme 3.2, there is a simplified flow diagram for the processing of a nanoclay. The production of a clay suitable for the preparation of polymer nanocomposites requires also a purification step to remove no-clay part inside the bentonite.

3.3 PURIFICATION OF CLAY Bentonite contains montmorillonite and other minerals. The overall composition depends on the geographical origin of the materials and on the mechanism they were formed through. Analysis such as X-ray diffraction (XRD) and the determination of the cationic exchange capacity (CEC) give information about the montmorillonite content inside a bentonite sample; the XRD of a clay taken directly from the mine shows typically bentonite together with feldspars, opal, quartz, calcite, and others. For applications such as the preparation of polymer nanocomposites, it is necessary a purification step of the mineral, with the aim to increase the montmorillonite content and to remove all the other minerals devoid of a lamellar structure. A good description of the methodology for clay separation is given by Bain and Smith [11]. The system is based on the different specific gravity of the minerals inside the bentonite. There are different levels of purification: it is first necessary to separate a <2 mm fraction by either sedimentation/decantation or centrifugation. Prior to separation, clay must be saturated with Na ions to provide a stable suspension and after separation excess salt must be removed from the clay product by sequential water and alcohol washing. When the clay has to be used for the preparation of polymer nanocomposites, the purification has to be particularly pushed, until bentonite consists essentially of montmorillonite. In the following, some details are given on the development of a specific purification process for nanoclay, which was pursued in the frame of a EU financed project. The purification process starts by preparing a clay suspension in water, allowing the heavier impurities to settle down. The suspension is then centrifugated in order to eliminate lighter impurities and, finally, the suspension is dried and grinded to the desired particle size. Design and optimization of the process conditions are done step by step, analyzing the material in the lab at each step, developing a specific process optimization for each clay.

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INDUSTRIAL TREATMENTS AND MODIFICATION OF CLAY MINERALS

FIGURE 3.2 Amount of material lost as a function of suspension concentration.

For a nanoclay production, the best choice is to select in the mine a raw material with a high level of purity, that is, that gives a mineralogical analysis corresponding to 100% smectite. 3.3.1 Influence of Clay Concentration The optimal clay concentration in the suspension has to be carefully determined. In fact, when the clay concentration is too high, the clay platelets cannot be evenly dispersed and the sedimentation step leads to a considerable loss of material, lowering the process productivity. Results from a systematic investigation on the amount of material lost as a function of suspension concentration are summarized in Figure 3.2. The highest mass loss shown in the bar chart was obtained when the suspension was decanted for 12 h and the supernatant was then centrifuged. A concentration of 7% (w/v) turned out to be the highest at which the mass loss was found not dependent on concentration, and was determined as the optimal value. 3.3.2 Influence of Swelling Time Time of clay swelling before suspension centrifugation is an important parameter, to be optimized. Suspensions with different time of swelling were prepared and analyzed with the same procedure adopted for concentration optimization; as shown in Figure 3.3, the optimal time is of 8 h, as shorter times leads to higher mass loss, while longer times seem not to give any advantage.

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FIGURE 3.3 Increasing swelling time decreases loss of mass in purification step.

3.3.3 Influence of Temperature The above studies led to the conclusion that for the considered clay the maximum concentration allowed for the clay suspension was 7% (w/v) and the minimum swelling time was 8 h; using these conditions to prepare the suspension, after the centrifugation there was a loss of 10% (loss of mass) of the starting material. The influence of temperature was then investigated preparing suspensions at temperatures from 35 to 65 C, fixing the swelling time at 8 h and the concentration at 7%, and checking the loss of mass to control the grade of purification. Figure 3.4 shows that at 50 C a good purification was obtained in these conditions.

FIGURE 3.4 Between 55 and 60 C the optimum temperature to disperse bentonite in water.

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INDUSTRIAL TREATMENTS AND MODIFICATION OF CLAY MINERALS

FIGURE 3.5 At fixed temperature (40 C) increasing swelling time decreases loss of weight in the purification step.

The trade-off between swelling temperature and swelling time was further investigated with two series of experiments at fixed temperature, reported in Figures 3.5 and 3.6. As can be easily seen doubling the swelling time (from 8 to 16 h) it is possible to prepare a good purified suspension also at 40 C. The data obtained at 50 C confirmed that 8 h is the minimum time to wait before the centrifugation process. Finally, the trade-off between swelling temperature and concentration was investigated with a last set of experiments at 60 C, reported in Figure 3.7: for a swelling time of 12 h, 9% seems to be a suitable concentration to use.

FIGURE 3.6 At fixed temperature (50 C) increasing swelling time decreases loss of weight in the purification step.

REACTION OF CLAY WITH ORGANIC SUBSTANCES

97

FIGURE 3.7 At fixed temperature and swelling time, increasing concentration increase loss of weight in purification step.

TABLE 3.2 Three Different Ways to Prepare the Suspension in Order to have the Same Grade of Purification

Suspension 1 Suspension 2 Suspension 3

Temperature ( C)

Swelling Time (h)

Concentration (w/v)

50 40 60

8 16 12

7 7 9

A suspension concentration of 9% means an increase of about 29% of the material that can be recovered at the end of the process and it is quite relevant for productivity. In summary, three different ways to prepare the suspension in order to have the same grade of purification were identified, as summarized in Table 3.2.

3.4 REACTION OF CLAY WITH ORGANIC SUBSTANCES The polarity of clay minerals can be tuned replacing, with organic cations, the inorganic metal ions saturating the structural negative charge on the silicate layers [12,13]. The reaction, shown in Figure 3.8, may be expressed as Q þ þ M þ -clay ! Q þ -clay þ M þ where Q þ is any organic cation.

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INDUSTRIAL TREATMENTS AND MODIFICATION OF CLAY MINERALS

Na+ N +

Na +2Cl–

Sodic bentonite

+

N

“QUAT”

N

+2NaCl

Organoclay

FIGURE 3.8 Reaction between ammonium salt and sodium–bentonite.

Organic cations might be quaternary onium salts (usually ammonium) or might be generated from an organic base capable of protonation in the aqueous solutions with subsequent pH adjustment to a suitable value. At a given pH, the concentration of cations in solution relative to the concentration of uncharged molecules is dependent on pK value. When cations are the dominant species in solution, the pH should be at least one or two units lower than the pK. If pH is too close to pK, cation exchange may be accompanied by adsorption of neutral molecules of the same organic species [14]. On the other hand, if the solution is too acidic, adsorption may be hindered due to competition with H þ ions or with metal cations released from silicate lattice by acid attack. Organic ion exchange will depend also on the solubility of the organic base in water, which, in turn, may be pH dependent. The accessibility of the exchange surface sites varies with the mineral system. Chemical and structural characteristics of the clay mineral have a marked influence on the rate of exchange. The total charge on the silicate layers and its location (octahedral and tetrahedral), as well as the size and valence of the saturating inorganic cations determine the strength of the electrostatic attraction between layers and influence the mobility of cations in interlayer space. Attraction between layers is greatest for higher charged silicates saturated with larger, less hydrated, 4þ monovalent cations, K þ , Rb þ , NH , Cs þ , leading to contraction of the basal  spacing approximately 9.8–10.8 A. On the other hand, with Na þ , Li þ , Ca2 þ , or Mg2 þ ions, because of their larger hydration energies, expansion to between 15 and 20 A or even to complete dissociation occurs, so that cations are more mobile and easier to replace. Water suspensions of Na–montmorillonite and of certain Li–vermiculite contain the mineral as fully dissociate layers, at large distance from each other. When organic cations are added to these suspensions all the clay surface is accessible to the solution phase, and exchange occurs readily. Most generally clay particles in suspension are domains consisting of stacks of elementary layers with d(001) spacing of less than  20 A, depending on the hydration properties of the interlayer cation. Exchange that is required by the organic cations to penetrate first into interlayer region to displace the inorganic metal ions attached to the internal surface. Virtually any organic cation  having less than 6 A as its smallest dimension can be introduced by ion exchange. The larger the size of the cation, the wider is the resulting spacing between adjacent layers in the complex formed.

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The reaction starts at the edges of the particle and proceeds toward the center in a highly regular fashion. The kinetics of the process has been studied by several authors [15,16]. Rate of penetration increases as the temperature of the reaction increases. Reaction temperatures in the 60–70 C range are recommended if they do not interfere with the stability of the organic material. Exchange can be completed after reaction times of the order of minutes for montomorillonites or hours for vermiculites. As exchange proceeds, the concentration in solution of the extracted metal cations may build up to critical equilibrium level. Montomorillonites are rather insensitive to their presence in solution and can adsorb organic cations even in presence of brines.

3.5 PARTICLE SIZE MODIFICATION A clay is “nano” in one dimension, the thickness, when it is fully dispersed in a matrix: water for inorganic clays, solvent for organoclays, polymer or rubber for nanoclay. The particle size of the material before dispersion, delamination, and exfoliation, is related to the application of the material. A clay finished product for many application could have a particle size from 75 mm or 45 mm as top size, on the other hand, typical nanoclay powder particle size must have 20 mm as top size. The reason is to facilitate the exfoliation of the nanoclay inside the matrix (polymer, rubber). So after the purification and surface modification, particle size also is an essenzial step of production process for nanoclay.

REFERENCES 1. Moore, D. M.; Reynolds Jr., R. C. (eds.), X-ray Diffraction and the Identification and Analysis of Clay Minerals, Oxford University Press, Oxford, 1997. 2. Christidis, G. E. Geological aspect an genesis of bentonite. Elements 5, 93–98 (2009). 3. Forano, C.; Hibino, T.; Leroux, F.; Taviot-Gu
eho, C. Layered double hydroxides. In Handbook of Clay Science, F. Bergaya, B. K. G. Theng, and G. Lagaly (eds.), Elsevier, Amsterdam, 2006, Vol. 1, pp. 1021–1095. 4. Beaufort, D.; Berger, G.; Lacharpagne, J.-C.; Meunier, A. An experimental alteration of montmorillonite to a di þ trioctahedral smectite assemblage at 100 and 200 C. Clay Min., 36, 211–225 (2001). 5. Knight, W. C. Bentonite. Eng. Miner. J. 66, 491 (1898). 6. Bergaya, F.; Lagaly, G.; Beneke, K. History of clay science: a young discipline. In Handbook of Clay Science, F. Bergaya, B. K. G. Theng, and G. Lagaly (eds.), Elsevier, Amsterdam, 2006, Vol. 1, pp. 1163–1181. 7. Grim, R. E.; G€uven, N. (eds.). Bentonites: Geology, Mineralogy and Uses, Elsevier, New York, 1978.

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8. Murray, H. H. Bentonite applications. In Applied Clay Mineralogy: Occurrences, Processing and Applications of Kaolins, Bentonites, Palygorskite-sepiolite, and Common Clays, H. H. Murray (ed.), Elsevier, Amsterdam, 2006, Vol. 2, pp. 111–130. 9. Lantenois, S.; Nedellec, Y.; Prelot, B.; Zajac, J.; Muller, F.; Douillard, J.-M. Thermodynamic assessment of the variation of the surface areas of two synthetic swelling clays during adsorption of water. J. Coll. interface Sci., 316(2), 1003–1011 (2008). 10. Christidis, G. E.; Blum, A. E.; Eberl, D. D. Influence of layer charge and charge distribution of smectites on the flow behaviour and swelling of bentonites. Appl. Clay Sci., 34, 125–138 (2006). 11. Bain, D. C.; Smith B. F. L. Chemical analysis. In A Handbook of Determinative Methods in Clay Mineralogy, M. J. Wilson (ed.), Blackie & Son Ltd., London, 1987, pp. 248–274. 12. Bergaya, F.; Lagaly, G. Clay minerals properties responsible for clay-based polymer nanocomposite (CPN) performance. In Clay-Based Polymer Nanocomposites. CMS Workshop Lecture, K. A.Carrado and F. Bergaya (eds.), The Clay Minerals Society, Chantilly, VA, 1997, Vol. 15, pp. 61–97. 13. Warr, L.; Berger, J. Hydration of bentonite in natural waters: application of “confined volume” wet-cell X-ray diffractometry. Phys. Chem. Earth, 32(1–7), 247–258 (2007). 14. Theng, B. K. G. Interactions of clay minerals with organic polymers. Some practical applications. Clays Clay Min., 18, 357–362 (1970). 15. Bain, D. C.; Smith B. F. L. Chemical analysis. In Handbook of Determinative Methods in Clay Mineralogy, M. Wilson (ed.), Chapman and Hall, New York, 1987, pp. 248–274. 16. Theng, B. K. G. The Chemistry of Clay—Organic Reactions, Adam Hilger Ltd, 1974, p. 343.

CHAPTER 4

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES HENDRIK HEINZ

4.1 STRUCTURE AND DYNAMICS The structure of alkylammonium chains on layered silicate surfaces has been first explored by Lagaly et al. [1–3] and refined by Vaia [4–6], Osman, Heinz, and other groups [7–35]. Understanding of structural features has been possible by X-ray diffraction, IR, AFM, NEXAFS, DSC, TGA, and TEM as well as by molecular simulation. Ultimately, accurate molecular models for mica-type silicates compatible with those for organic molecules, polymers, and biopolymers [14,24] have enabled the visualization of alkylammonium-modified silicates and computation of structural, thermal, surface, and mechanical properties in quantitative agreement with measurements [14,16,18,24,27,29,32–34]. A wide range of interlayer structures and cohesive forces of alkylammonium surfactants on layered silicate surfaces is found as a consequence of the polarity of the mineral surface and the distribution of local charge defects in the tetrahedral and octahedral layers. These charge defects can be described as SiO2 ! AlO2 . . .Kþ and AlO(OH) ! AlO(OH) . . . Naþ substitution and cause the presence of interlayer cations [36]. The properties of organically modified clay minerals then depend on the cation exchange capacity (CEC) and on the surfactant structure, that is, nature of the head group, number of chains per head group, and chain length. The multipolar environment created by the aluminosilicate surface is best described by atomic charges of about þ1.1e for tetrahedral Si, þ1.45e for octahedral Al, 0.55e for tetrahedral O, and   0.76e for octahedral O atoms. At the defect sites, the charge of tetrahedral Al and octahedral Mg is somewhat lower (þ0.8e and þ1.1e) and the atomic charges of surrounding oxygen atoms are increased (  0.78e and   0.87e) to accommodate the negative charge that balances the positives charge of interlayer cations [4,36].

Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

4.1.1 Packing Density and Self-Assembly The range of possible homogeneous alkyl structures on layered silicate surfaces is schematically depicted in Figure 4.1 [32]. The packing density l0 is a major order parameter and represents the average cross-sectional area per chain AC;0 in relation to the available surface area AS per chain: l0 ¼

AC;0 AS

ð4:1Þ

The packing density can range between 0 and 1 whereby a low packing density corresponds to a large amount of surface area per chain and a high packing density corresponds to small surface area per chain close to the cross-sectional area of the chain. This concept is practical because the surface area per alkyl chain, in all-anti conformation, is a constant equal to AC;0 ¼ 0:188 nm2, and the surface area per alkyl chain AS is given by the CEC, for example, AS ¼ 0:468 nm2 for mica. The packing density is directly related to the average segmental tilt angle u0 of the alkyl chains on the layered silicate surface relative to the surface normal: l0 ¼ cos u0

ð4:2Þ

Visualizations of the surface structure of alkylammonium montmorillonites as a function of CEC, length of the alkyl chains, and head group structure in all-atomic detail are

FIGURE 4.1 Structure and dynamics of homogeneous alkyl layers (chain length >C10) grafted to even clay mineral surfaces. (a) At low packing density, we find disordered alkyl chains oriented parallel to the surface. (b) At intermediate packing density, alkyl chains are of intermediate order and of an intermediate collective tilt angle relative to the surface. (c) At high packing density, the surfactant chains are nearly vertically oriented and all-anti configured. The packing density is directly proportional to the cation exchange capacity and the number of alkyl chains per surfactant head group. Reversible thermal transitions are observed for intermediate packing density (in (b)). (Reprinted from Ref. [32] with permission by the American Chemical Society.)

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FIGURE 4.2 Representative snapshots of montmorillonite of low CEC (91 mequiv./100 g) with trimethylalkylammonium surfactants N(CH3)3þ  CnH2n þ 1 at equilibrium distance. The formation of partially and densely packed alkyl layers can be seen. The quaternary ammonium head group cannot form hydrogen bonds to the surface and remains shared between the two clay layers at equilibrium distance until a partial alkyl bilayer is formed. (Reprinted from Ref. [35] with permission by Taylor and Francis.)

shown in Figures 4.2–4.4. We include a montmorillonite of low CEC (91 mequiv./ 100 g) with the approximate formula Na0.333[Si4O8][Al1.667Mg0.333O2(OH)2], a montmorillonite of high CEC (143 mequiv./100 g) with the approximate formula Na0.533[Si4O8][Al1.467Mg0.533O2(OH)2], and muscovite mica of very high CEC (251 mequiv./100 g) with the formula K[Si3AlO8][Al2O2(OH)2]. These examples correspond to a smectite, a vermiculite, and a mica with an increasing packing density of monoalkylammonium chains in the order 0.13, 0.21, and 0.40. Dialkylammonium chains C12 and C18 on mica in Figure 4.4 represent a packing density of 0.8. Each of these cases is associated with particular structural and dynamic features, which we describe in the following sections. At a low CEC of 91 mequiv./100 g (Figure 4.2), the progression of the basal plane spacing in XRD occurs in a plateau-and-step fashion with increasing chain length as shown in Figure 4.5a [29]. This characteristic pattern is related to layer-by-layer assembly of the alkyl chains. Up to a chain length C8 with NMe3 head groups

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ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

FIGURE 4.3 Representative snapshots of montmorillonite of high CEC (143 mequiv./100 g) with alkylammonium surfactants H3N þ  CnH2n þ 1 at equilibrium distance. The formation of partially packed and densely packed alkyl monolayers, bilayers, and pseudotrilayers can be seen. Head groups are positioned very close to the surface due to hydrogen bonds and, therefore, are vertically separated between the two clay layers at equilibrium distance even for short-chain length. (Reprinted from Ref. [35] with permission by Taylor and Francis.)

(Figure 4.2), a densely packed alkyl monolayer is fitted into the interlayer space. Beginning at chain length C10, the interlayer gallery is pushed apart by partial formation of an alkyl bilayer (Figures 4.2 and 4.5a). This alkyl bilayer is assembled at a low interlayer density at a chain length C14 and densely packed at a chain length C18 (Figure 4.2). At a chain length C22, another upward push in the gallery spacing occurs due to the formation of a partial alkyl trilayer (Figures 4.2 and 4.5a). The subsequent formation of partially and densely packed alkyl layers can be monitored more precisely by the interlayer density (Figure 4.5b), that is, the density of organic material in the interlayer space excluding the volume of the unmodified aluminosilicate layer. The thickness of the unmodified aluminosilicate layer corresponds to the basal plane spacing of pyrophyllite of 0.919 nm, equal to mica and montmorillonite without interlayer cations [29,37]. The interlayer density for the series with

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105

FIGURE 4.4 Representative snapshots of mica of very high CEC (251 mequiv./100 g) with trimethyloctadecylammonium surfactants (CH3)3Nþ  C18H37, dimethyldidodecylammonium surfactants (CH3)2N þ (C12H25)2, and dimethyldioctadecylammonium surfactants (CH3)2N þ (C18H37)2 at equilibrium distance at 0 C. The increase in basal plane spacing and the decrease in tilt angle relative to the surface normal can be seen for the single alkyl surfactant in comparison to the double alkyl surfactant. The nominal packing density doubles from 0.40 for C18-mica to 0.80 for 2C12-mica and 2C18-mica. (Reprinted from Ref. [14] with permission by the American Chemical Society and from Ref. [16] with permission by Wiley-VCH.)

NMe3 head groups indicates a steady increase in gallery spacing up to chain length C8 that corresponds to a densely packed alkyl monolayer (Figure 4.2). For chain lengths C10 and C12, the interlayer density drops due to formation of a partial alkyl bilayer and then increases up to chain length C18/C20 as the alkyl bilayer experiences denser packing. A drop in interlayer density occurs again at chain length C22 due to formation of a partial alkyl trilayer that cannot fit into the space of the alkyl bilayer. These fluctuations in interlayer density have implications on the cleavage energy of these alkylammonium montmorillonites, which we will explain in Section 4.3. Also, fluctuations in interlayer density are maximal for the organically modified clay minerals in the dry state and mitigated under ambient conditions due to the attraction of moisture, which can fill the interlayer space in case of otherwise low interlayer density. The variable interlayer density also imposes conformational constraints on the alkyl backbones. The degree of confinement of the alkyl tails between the aluminosilicate layers is proportional to the percentage of gauche conformations (Figure 4.5c). The preferred conformation of the alkyl surfactants at low temperature is all-anti that is approximated with 15% gauche conformation for loosely packed alkyl chains in the presence of significant available interlayer space (C4 and C16 in Figure 4.5c). Energetically more demanding conformations of 35% gauche content are enforced for densely packed alkyl chains due to shortage of available interlayer space (Figure 4.5c). High interlayer density thus necessitates increased chain folding to avoid further expansion of the interlayer that would lead to a more significant loss

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

(a)

50 NH3 head group

Percentage gauche conformation

45

NMe3 head group

40 35 30 25 20 15 10 0

2

4

6

8 10 12 14 16 18 Carbon atoms per chain

20

22

24

(b)

Interlayer density (kg/m1)

800

700

600

500 400 NH3 head group NMe3 head group

300 0

2

4

6

8

10

12

14

16

18

20

22

24

Carbon atoms per chain

(c) 23 22 21 Basal plane spacing (A)

106

20 19 18 17 16 15 14

NH3 head group

13

NMe3 head group

12

0

2

4

6

8

10

12

14

16

Carbon atoms per chain

18

20

22

24

STRUCTURE AND DYNAMICS

107

in internal energy. This trend is clear for surfactants with NMe3 head group (Figure 4.5c) and altered for surfactants with NH3 head groups due to the formation of hydrogen bonds with the montmorillonite surface (Figure 4.3). The hydrogen bonds enforce gauche conformations near the head group, which significantly increase the percentage of gauche conformations, especially for short chains (Figure 4.5c). This bias in percentage of gauche conformations decreases toward longer chains with a higher total number of backbone dihedral angles. The fundamental trend in backbone conformation for hydrogen-bonded primary ammonium surfactants is the same as for quaternary ammonium surfactants, after subtracting the superimposed effect of the hydrogen-bonded head group. At a medium CEC of 143 mequiv./100 g (Figure 4.3), the progression of the basal plane spacing in XRD with increasing chain length is more steadily compared to lower CEC as shown in Figure 4.6a [29]. The curve still contains plateaus and steps, however, they are much weaker than observed at low CEC as the layer-by-layer assembly of the alkyl chains transitions toward the growth of uniform alkyl monolayers with a tilt angle significantly below 90 relative to the surface normal (see Eqs. (4.1) and (4.2)). For the series of NH3 head groups (Figures 4.3 and 4.6a), a densely packed “flat-on” alkyl monolayer is formed at chain length C6, a densely packed alkyl bilayer at chain length C14, and a densely packed alkyl trilayer at chain length C20. Nevertheless, these layers appear more disordered than at low CEC and are effectively tilted slightly upwards with respect to the surface (Figure 4.3). Thus, the gallery spacing is also higher than expected for flat-on alkyl monolayers, bilayers, and trilayers (Figure 4.6a). The interlayer density indicates minima and maxima for partially and densely packed alkyl layers (Figure 4.6b) although these fluctuations converge to a value of 725 kg/m3 for alkyl multilayers. Even the alkyl bilayer is difficult to distinguish from the alkyl trilayer taking into account the uncertainty (Figure 4.6b). Similarly, fluctuations in average chain conformation (Figure 4.6c) are significantly decreased compared to low CEC (Figure 4.5c) and converge to a percentage of 25–30% gauche conformations at 298 K for alkyl chains longer than C10. For comparison, the percentage of gauche conformations in liquid octadecylamine at 100  C is 39% [24]. Alkylammonium chains grafted to montmorillonite at high CEC are thus in a liquid-like state. The influence of hydrogen bonds of the primary ammonium head groups to the montmorillonite surface leads to a peak in the percentage of gauche conformations for the short C6 surfactant (Figures 4.3 and 4.6c) that has also been observed by IR spectroscopy [4,29] At high CEC of 251 mequiv./100 g in mica (Figure 4.4), the progression of the basal plane spacing in XRD with increasing chain length is linear due to the formation 3 FIGURE 4.5 (a) Computed basal plane spacing (0–5% deviation from XRD measurement), (b) interlayer density, and (c) average conformation of the alkyl chains between the silicate layers for alkylammonium-modified montmorillonite with CEC ¼ 91 mequiv./100 g. The series NH3þ –CnH2n þ 1 and N(CH3)3þ –CnH2n þ 1 with n ¼ 2, 4, . . ., 22 are shown. (Reprinted from Ref. [29] with permission by the American Chemical Society.)

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

(c)

50

NH3 head group NMe3 head group

Percentage gauche conformation

45 40 35 30 25 20 15 10 0

(b)

2

4

6

8 10 12 14 16 18 Carbon atoms per chain

20

22

24

1000 950

Interlayer density (kg/m1)

900 850 800 750 700 650 600 550

NH3 head group NMe3 head group

500 450

0

2

4

6

8 10 12 14 16 18 Carbon atoms per chain

20

22

24

(a) 30 28 26 Basal plane spacing (A)

108

24 22 20 18 16

NH3 head group NMe3 head group

14 12 0

2

4

6

8 10 12 14 16 Carbon atoms per chain

18

20

22

24

STRUCTURE AND DYNAMICS

109

FIGURE 4.7 Formation of (a) a spatially homogeneous phase and (b) two spatially separated phases of alkali ions and alkylammonium ions on the surface of layered silicates. The basal plane spacing d of the two-phase structure is higher than that of the corresponding one-phase structure. (c) Observation of spatially separated phases by contact-mode AFM on a mica surface coated with 0.1 mM trimethyloctadecylammonium bromide solution (B) after 24 h and (F) after 168 h. (Reprinted from Ref. [8] with permission by the American Chemical Society.)

of tilted oriented alkyl monolayers (see Figure 4.1b and c) [32]. The interlayer density is essentially constant as a function of chain length and the percentage of gauche conformations in the alkyl backbones is 20% for C18 chains at a packing density of 0.4, and only 10–15% for 2C12 and 2C18 chains at a packing density of 0.8 at temperatures between 0 and 20 C. The high packing density also allows spatial segregation of cations and alkylammonium surfactants when exchange of superficial alkali ions for alkylammonium ions is less than quantitative. This effect has been observed by AFM [8] and analyzed by a combination of XRD and molecular models. Spatially homogeneous structures with alkylammonium surfactants and interspersed remaining alkali cations are formed on the surface, or spatially separated islands of alkylammonium surfactants as well as regions containing alkali ions only, see Figure 4.7 [16]. The preference for a given structure as a function of grafting density and chain length can be predicted on the basis of enthalpic and entropic considerations [16]. 3 FIGURE 4.6 (a) Computed basal plane spacing (0–5% deviation from XRD measurement), (b) interlayer density, and (c) average conformation of the alkyl chains between the silicate layers for alkylammonium-modified montmorillonite with CEC ¼ 143 mequiv./100 g. The series NH3þ –CnH2n þ 1 and N(CH3)3þ –CnH2n þ 1 with n ¼ 2, 4, . . ., 22 are shown. (Reprinted from Ref. [29] with permission by the American Chemical Society.)

110

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

FIGURE 4.8 The organic–inorganic interface. (a) R–NH3þ head groups form hydrogen bonds to oxygen on the silicate surface. The average O    H distance is 150 pm and the mean N to O distance is 245 pm. (b) R–N(CH3)3þ head groups do not form hydrogen bonds to the surface. The average O    H distance is almost twice as large, 290 pm, and the mean N to O distance is 390 pm. Qualitatively, this implies the mobility of the R–NMe3þ on the surface should be higher than that of R–NH3þ . (Reprinted from Ref. [29] with permission by the American Chemical Society.)

4.1.2 Dynamics and Diffusion at the Clay–Surfactant Interface The interface between the clay mineral and alkylammonium surfactants is characterized by ionic bonds due to the exchange of interlayer alkali cations and can involve additional hydrogen bonds, particularly in primary alkylammonium surfactants, as shown in Figure 4.8 [29]. Hydrogen bonds decrease the lateral mobility of head groups and superficial diffusion. While the height of barriers for lateral movement of surfactant head groups varies, alkylammonium surfactants are mobile on the layered silicate surfaces at any given CEC. The diffusion mechanism involves the displacement of a surfactant head group to another cavity on the aluminosilicate surface and the concomitant movement of the head group of another surfactant, or a group of surfactants, to cover the then-empty cavity on the aluminosilicate surface. The correlation of the motion between a group of at least two surfactants results from the desire of the positively charged head groups to remain close to the spatially fixed negative charge defects inside the aluminosilicate layer. The fastest movement occurs for short quaternary ammonium surfactants on the outer surface of the modified clay surface. 2D self-diffusion constants of 5:0  106 cm2/s for Me3NEtþ and for Me3NBuþ were computed at CEC ¼ 91 mequiv./100 g at room temperature, which amounts to 20 and 6% of the 3D self-diffusion constant of liquid water (2:3  105 cm2/s). Self-diffusion constants for hydrogen bonded surfactants, surfactants at higher CEC, and surfactants in twolayer structures are <107 cm2/s. An important consequence of the ease of movement of quaternary ammonium surfactants in particular is the reduction of the diffusion barrier at higher temperature. For longer quaternary ammonium surfactants, the activation of lateral rearrangement leads to a second thermal transition observable in DSC in addition to the reversible melting transition of the alkyl backbone of the surfactant (see Section 4.2.1) [14,16].

THERMAL PROPERTIES

111

4.1.3 Utility of Molecular Simulation to Obtain Molecular-Level Insight Molecular dynamics simulation on the basis of interatomic potentials (force fields) and Monte Carlo methods on the basis of coarse-grain models have become useful tools to understand the molecular structure of alkylammonium-modified layered clay minerals. Fully atomistic molecular dynamics simulations of several 10,000 atoms can reach millisecond time scales and coarse-grain simulations corresponding to millions of atoms can reach beyond second time scales using state-of-the-art computational resources. Qualitative predictions have been possible early on [7,11,13,15,19,23,25,26], and the improvement of force fields models for quantitative simulation of polarity, surface, and interface properties [14,16,18,24,27] enables the simulation of basal plane spacing in 0–5% agreement with experiment, tilt angles in quantitative agreement with NEXAFS data, conformational data in quantitative agreement with IR and NMR data, and surface tensions as well as cleavage energies in 0–10% agreement with laboratory measurements [29,32–35]. For the computation of mechanical properties on the nanoscale, density functional theory (DFT) for small unit cells of the minerals yields results in even better agreement with experiment compared to classical force fields. Current force fields overestimate in-plane elastic moduli of the clay layers due to limitations in the choice of classical energy terms [37] while interlayer and shear properties between silicate layers can be computed reliably (Section 4.4).

4.2 THERMAL PROPERTIES Thermal properties of organoclays play an important role in processing and stability of nanocomposites. The presence of hydrocarbon-based surfactants can lead to reversible melting transitions upon heating to temperatures below 100 C, the possibility of moisture occlusion in the interlayer can lead to evaporation of adsorbed interlayer water upon first heating above 100 C, and thermal decomposition of alkylammonium surfactants proceeds above 200 C.

4.2.1 Reversible Melting Transitions of Alkyl Chains in the Interlayer Phase transitions at temperatures below 100 C are observed in alkylammoniummodified clay minerals of intermediate packing density between 0.20 and 0.75 (Figure 4.1). This packing density and thermal transitions are found in montmorillonites of low cation exchange capacity modified with dialkylammonium, trialkylammonium, and tetraalkylammonium surfactants, and in mica modified with monoalkylammonium surfactants. The phase transitions have been observed by DSC as shown in Figure 4.9 and are accompanied by an increase in basal plane spacing on the order of 5%, changes in IR frequencies of the symmetric and of the antisymmetric CH2 stretching vibration (ns;CH2 and na;CH2 ) by a few cm1, and changes in the 13C NMR spectral shifts of backbone CH2 groups [4,10,12,17,22]. Notably, two-phase transitions can be observed for quaternary alkylammonium

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

112

59 69ºC

2924

57

2922 52ºC

53 51 49

vmax (cm–1)

Heat flow (mW)

43ºC

55

ν25 CH2str

2920 2918 2916

39ºC

47 –30 –10 10 30 50 70 90 110 130 Temperature (ºC) (a)

2914 20

80 60 40 Temperature (ºC) (b)

FIGURE 4.9 (a) DSC trace of octadecyltrimethylammonium mica (80% ion exchange and packing density 0.40) that shows two transitions upon heating of which the second transition at 69 C is not immediately reversible. The insets show the corresponding molecular structure at 20 and 100 C as seen in molecular dynamics simulation. (b) The corresponding change in the asymmetric CH2 stretching vibration. Infrared spectroscopy only monitors major conformational changes during the second transition near 69 C. (Reprinted from Refs [10] and [14] with permission by the American Chemical Society.)

surfactants at medium packing density such as in N(CH3)3C18H37-mica (Figure 4.4). The transition at lower temperature represents a partial order–disorder (melting) transition of the rod-like surfactants into disordered rods. The transition at higher temperature involves lateral displacements of the quaternary ammonium head groups on the mica surface, which facilitates a higher degree of disorder of the alkyl backbones and involves a more substantial melting enthalpy (12 kJ/mol) [10,14]. Lateral displacement of the head groups has been observed by molecular dynamics simulation (Figure 4.10) [14], and the second transition at higher temperature is only recovered after cooling for several hours because passing the substantial activation barrier for reverse lateral displacement of surfactant head groups requires more time at lower temperature [10,14]. Thermal phase transitions have also been reported experimentally for 2Cn, 3Cn, and 4Cn surfactants on montmorillonite of low CEC (72 mequiv./100 g) by DSC, IR, NMR, and XRD [17]. The transitions of three-armed and four-armed surfactants (packing densities 0.33 and 0.44) are of more substantial melting enthalpy in comparison to the two-armed surfactants (packing density 0.22). The use of primary ammonium surfactants instead of quaternary ammonium surfactants leads to higher barriers for lateral rearrangement due to hydrogen bonds to the surface (Figure 4.8) and results in only one reversible melting transition

THERMAL PROPERTIES

113

FIGURE 4.10 Distribution of the trimethylammonium head groups of N(CH3)3(C18H37)þ cations on the mica surface, represented by the green nitrogen atoms at (a) 20 C and (b) 100 C after 400 ps of MD simulation. With increasing temperature, the cationic surfactant head groups move laterally across the cavities in the mica surface. (Reprinted from Ref. [14] with permission by the American Chemical Society.)

upon heating. This behavior is shown in Figure 4.11a for dioctadecylammonium surfactants on montmorillonite of CEC 72 mequiv./100 g. 4.2.2 Solvent Evaporation and Thermal Elimination of Alkyl Surfactants Evaporation of residual water proceeds at temperatures above 100 C. The hydrophilicity of clay mineral surfaces and, particularly, the formation of loosely packed interlayer structures (Section 4.1.1) at low packing density lead to the inclusion of water or other solvent molecules in the interlayer space. Upon first heating above the boiling point of the solvent, residual solvent evaporates and the weight loss can be registered by DSC (Figure 4.11a). The effect typically disappears upon second heating. Thermal elimination of alkyl surfactants proceeds at temperatures above 200 C and is illustrated for the example of alkylammonium ions. The thermal decomposition of tetraalkylammonium micas and tetraalkylammonium montmorillonites proceeds analogous to the cleavage of quaternary ammonium hydroxides by Hofmann elimination with an E2 mechanism. In this mechanism, the mica-multianion and the montmorillonite-multianion act as a medium strong base with the corresponding acids H-mica and H-montmorillonite after elimination [14]. mont     NR3þCH2 CH2 R0

! mont    H 

þ

þ NR3 þ CH2 ¼ CHR0 ð4:3Þ

The alkylammonium surfactants sterically achieve the antiperiplanar transition state for b-elimination by alignment of the first two carbon atoms departing from the

114

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

(a)

–1

Heat flow [mW]

1.5

ng

ati

2C18

e .H

Loss of solvent

1

ating

2. He

0.5

ng

–0.5

Cooli

–1.5 –30

–10

10

30

50

70

90

100

130

150

Temperature (ºC) (b)

dm/dt

4C18 3C18

2C18 C18 100

200

300

400

500

600

700

800

900

Temperature (ºC)

FIGURE 4.11 (a) DSC trace of dioctadecylammonium-modified montmorillonite (CEC 72 mequiv./100 g). A minor reversible melting transition at 55 C and loss of solvent (EtOH/ H2O) upon first heating at 75  140 C is seen. The packing density of 0.22 is just above the threshold of 0.20 for the occurrence of reversible thermal transitions (Figure 4.1). (b) A DTG plot for multiarmed surfactants on montmorillonite (CEC 72 mequiv./100 g). Thermal decomposition according to Eq. (4.3) begins in the temperature range 200240 C and reaches a maximum at 280320 C. (Reprinted from Ref. [17] with permission by the American Chemical Society.)

nitrogen atom roughly parallel to the surface. Especially at higher temperature, conformational flexibility supports arrangements close to the requirements of the transition state of the E2 elimination. The b-hydrogens on the ammonium ions approach the level of the upper oxygen atoms of the mica surface about 300 pm and H-abstraction can occur. The usual elimination temperature for quaternary hydroxides is 100–200 C. Mica is a weaker base compared to hydroxide ions and, therefore, the elimination temperature is roughly 250 C. Further evidence can be obtained by a detailed analysis of the elimination products.

LAYER SEPARATION AND MISCIBILITY WITH POLYMERS

115

From a practical viewpoint, thermal instability of alkylammonium montmorillonites is an undesirable effect during melt processing of clay–polymer composites since temperatures are often above 200 C. A possible approach to increase thermal stability up to higher temperatures may be the usage of alkylammonium surfactants substituted in the b-position NR3þ –CH2–CR1R2–R0 . Also, additional substitution of H atoms in the a position can be invoked. Then, no hydrogen atoms will be available for the elimination reaction and future research will yet be necessary to show the benefits and limitations of such approaches.

4.3 LAYER SEPARATION AND MISCIBILITY WITH POLYMERS The challenge of dispersion of organically modified clay minerals in polymer matrices is directly connected with their surface and interface properties. Quantitative understanding of these properties relies on the detailed knowledge of the structure and dynamics established in the previous sections. 4.3.1 Thermodynamics Model for Exfoliation in Polymer Matrices During dispersion in a polymer matrix, the minerals (M) give up their cohesive free energy DGM that equals the free energy of cleavage, the polymer (P) gives up locally its cohesive free energy per area DGP , and the two components form a new interface (MP) associated with a free energy DGMP [34]. The process is schematically depicted in Figure 4.12 and the overall change in free energy DG is: DG ¼ DGM þ DGP þ DGMP

ð4:4Þ

The first two terms are positive, the third negative. The advantage of this free energy model is the focus on dispersion aspects and the full inclusion of entropic effects of the polymer in the second and third term [34]. Under the assumption that all surfaces would be well defined and do not reconstruct, Eq. (4.4) can be rewritten per interfacial area using surface and interface tensions: gMP ¼ gM þ gP þ DGMP

ð4:5Þ

ΔG

Polymer

FIGURE 4.12 Schematic representation of exfoliation of mineral layers (M) in a polymer matrix (P) and the associated free energy change. (Reprinted from Ref. [34] with permission by the American Chemical Society.)

116

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

gMP reflects the definition of the mineral–polymer interface tension and Eq. (4.5) can be imported into Eq. (4.4) to yield DG ¼ DGM  gM þ DGP  gP þ gMP

ð4:6Þ

Equation (4.6) states that the free energy of exfoliation can be lowered not only by reduction of the interface tension gMP but also by cleavage free energies lower than the corresponding surface tensions of the mineral and of the polymer. We will show in the following that the formation of partially and densely packed alkyl layers leads to significant variation in cleavage energy in a range from 25 to 200 mJ/m2 while corresponding surface tensions of the cleaved mineral layers remain in a narrow range of 40–45 mJ/m2 [21,38–40]. Therefore, modified clay minerals of low cleavage energy are expected to better disperse in polymer matrices than those of higher cleavage energy. Very high cleavage energies in the range 100–375 mJ/m2 are also found for the natural (unmodified) clay minerals as well as for certain organically modified montmorillonites due to strong Coulomb contributions. In the absence of significant Coulomb contributions due to design of head groups and suitable surfactants, cleavage energies of organically modified montmorillonites can be significantly lower than the surface tension and enhance exfoliation according to Eq. (4.6). 4.3.2 Cleavage Energy The cleavage free energy and the surface tension (Eq. (4.6)) differ due to surface reconstruction during cleavage. DGM may vary between slow (equilibrium) and fast (nonequilibrium) cleavage, it can be very high when interlayer cations are separated upon cleavage, it can be low for loosely packed alkyl chains in the interlayer, and it differs for densely packed versus loosely packed alkyl chains between the clay mineral layers. The cleavage free energy differs from the surface tension that is determined for surfaces after cleavage in contact with different liquids [27]. The cleavage free energy and the surface may converge when surfaces undergo minimal reconstruction upon cleavage or when reconstruction upon cleavage is similar to reconstruction upon interaction of the cleaved surface with liquids during contact angle measurements. We define the cleavage energy eS (focusing on the energy contribution only) per contact surface area 2A as the difference in average total energy between a cleaved assembly of layers at infinite separation DES and the united assembly of layers at equilibrium distance DEU, assuming a slow cleavage process in thermodynamic equilibrium, as illustrated in Figure 4.13: eS ¼

DES  DEU 2A

ð4:7Þ

The analysis of cleavage energies extends the structural analysis in Section 4.1.1. The concepts of gallery spacing, interlayer density, and chain conformation allow a

LAYER SEPARATION AND MISCIBILITY WITH POLYMERS

117

FIGURE 4.13 Illustration of the cleavage process for montmorillonite (CEC 143 mquiv./ 100 g) modified with tetraethylammonium ions. The cleavage energy equals the difference in total energy between (a) the separated structure of energy ES and (b) the unified structure of energy EU per contact surface area 2A. During the cleavage process, vertical separation of the head groups between the two montmorillonite layers, lateral movements of the head groups, and conformational reconstruction of the alkyl backbone occur. (Reprinted from Ref. [35] with permission by Taylor and Francis.)

systematic understanding of cleavage energies as a function of CEC, head group chemistry, and chain length. The cleavage energies are presented in Figure 4.14 for alkylammonium montmorillonites of CEC 91 meq/100 and 143 mequiv./100 g with surfactants containing NH3 and NMe3 head groups (see Figures 4.2 and 4.3). We observe a higher cleavage energy of short quaternary alkylammonium surfactants and consistently lower cleavage energy of primary alkylammonium surfactants. The distinction arises from differences in the vertical distribution of the cationic head groups between the layers at equilibrium separation. Bulky quaternary ammonium head groups are located in the center of the interlayer for short-chain lengths (Figure 4.2) that leads to strongly attractive Coulomb interactions between each quaternary cation and anionic defect sites in both clay layers, as well as some opposing repulsion by cation neighbors in the interlayer. Similar to unmodified clay minerals [27], the cleavage energy is then very high and decreases only after a partial second alkyl layer causes the head groups to vertically separate between the layers, thereby breaking the strong Coulomb attraction across the interlayer (C14 in Figure 4.2). Thus, the cleavage energy of montmorillonite with quaternary alkylammonium ions decreases at a chain length equivalent to a partial alkyl bilayer (Figure 4.14). In the case of primary alkylammonium ions, head groups are vertically separated between the two layers even for the shortest surfactants due to their small

118

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

(a) Cleavage energy per contact area (mJ/m2)

160 NMe3 head group

140

NH3 head group

1

CEC— 91 mequiv./100 g

120 100 80

2 2

60

1

40 20 2

4

6

8 10 12 14 16 18 Carbon atoms per backbone

20

24

22

Cleavage energy per contact area (mJ/m2)

(b) 220

NMe3 head group

200

1

NH3 head group

180

CEC—91 mequiv./100 g

160 140 120 100 80 60

2

40

2

3

3

4

1

20 2

4

6

8 10 12 14 16 18 Carbon atoms per backbone

20

22

24

FIGURE 4.14 Cleavage energy of two montmorillonite layers modified with Cn-alkylammonium ions as a function of chain length n and head group structure (NH3 and NMe3). (a) Low CEC (91 mequiv./100 g) and (b) high CEC (143 mequiv./100 g). The labels 1, 2, 3, 4 indicate the chain length at which a densely packed alkyl monolayer, bilayer, pseudotrilayer, and pseudoquadrilayer are formed in the interlayer space. (Reprinted from Ref. [34] with permission by the American Chemical Society.)

LAYER SEPARATION AND MISCIBILITY WITH POLYMERS

119

size and strong binding to the surface of only one clay layer supported by hydrogen bonds (C2 in Figure 4.3). Thus, Coulomb interactions across the interlayer are then so much reduced that the cleavage energy is mostly determined by van der Waals interactions between the short alkyl chains. The cleavage energy changes periodically with the formation of densely packed and loosely packed alkyl layers, and only minor Coulomb contributions are involved (Figures 4.14 and 4.15). In particular, the distinctively low cleavage energies for montmorillonite modified with short primary ammonium ions may support exfoliation in polymer matrices by lowering DGM (Eq. (4.6)). We will explain the relationships between interlayer structure and cleavage energy in detail for organically modified montmorillonites of low CEC (91 mequiv./100 g), including contributions by Coulomb and van der Waals energy to the cleavage energy (Figure 4.15). For surfactants with NH3 head groups (Figure 4.15a), the cleavage energy follows a trend closely associated with packing of the alkyl layers in the interlayer space. Lower cleavage energy is observed in loosely packed alkyl layers since less van der Waals interactions are removed across the interlayer space upon cleavage. Higher cleavage energy is observed in densely packed alkyl layers as more van der Waals interactions are removed across the interlayer space upon cleavage resulting. Lowest cleavage energies of 25 and 33 mJ/m2 are computed for C4 and C14 surfactants, corresponding to a loosely packed alkyl monolayer and a loosely packed alkyl bilayer in the interlayer space, respectively. Highest cleavage energies of 46 and 52 mJ/m2 are seen for C10 (C8) and C22 surfactants, corresponding to a densely packed alkyl monolayer and a densely packed alkyl bilayer in the interlayer space, respectively. A slightly increased value for C2 surfactants arises from residual Coulomb interactions. The most remarkable difference to quaternary ammonium head groups is the small magnitude of Coulomb interactions <20 mJ/m2 due to the vertical distribution of head groups between the two montmorillonite layers at equilibrium separation. The attachment of the small primary, hydrogen-bonded head groups to either one of the two montmorillonite layers at equilibrium separation can be seen, at CEC 143 mequiv./100 g, in Figure 4.3 Residual Coulomb energy decreases with increasing chain length to minor values beyond C8 (Figure 4.15a). The van der Waals energy increases steadily toward a densely packed alkyl monolayer at C10, followed by a decrease for a partial alkyl bilayer and another increase toward a densely packed alkyl bilayer at C22. The trend in cleavage energy follows approximately the trend in van der Waals contributions and the uncertainty is 3 mJ/m2. For surfactants with NMe3 head groups (Figure 4.15b), the cleavage energy reaches a maximum near 126 mJ/m2 for the chain length C4 and C6, goes through a minimum near 42 mJ/m2 at C12, reaches a second maximum of 65 mJ/m2 at C20, and falls to 39 mJ/m2 at C22. High cleavage energies for chain lengths up to a partial alkyl bilayer (C10) result from significant Coulomb contributions up to 80 mJ/m2 associated with vertical sharing of head groups between the two montmorillonite layers at equilibrium separation. The head groups separate vertically between the two layers upon formation of a partial alkyl bilayer in the interlayer space at chain length C12, leading to a significant decrease in Coulomb energy. The difference in head group

120

ALKYLAMMONIUM CHAINS ON LAYERED CLAY MINERAL SURFACES

Cleavage energy per contact area (mJ/m2)

(a) 120

Total VoW Coulomb Internal

NH3 head group CEC—91 mequiv./100g

100 80 1

2

60 40 20 0 –20 2

4

6 8 10 12 14 16 18 20 22 Carbon atoms per surfactant backbone

24

Cleavage energy per contact area (mJ/m2)

(b) 160

NMe3 head group

140

CEC—91 mequiv./100 g

120

Total VoW Coulomb Internal

100 80 60 40 1

20

2

0 –20 2

4

8 10 12 14 16 18 20 22 6 Carbon atoms per surfactant backbone

24

FIGURE 4.15 Cleavage energy and additive contributions from Coulomb, van der Waals, and internal interactions for montmorillonite of low CEC (91 mequiv./100 g) modified with (a) alkylammonium ions with NH3 head groups and (b) alkylammonium ions with NMe3 head groups. The vertical lines designated 1 and 2 indicate the chain length corresponding to a densely packed flat-on alkyl monolayer and a densely packed flat-on alkyl bilayer in the interlayer space. (Reprinted from Ref. [34] with permission by the American Chemical Society.)

distribution is illustrated in Figure 4.2 for C2 (no vertical separation at equilibrium distance) and C14 (vertical separation at equilibrium distance). The van der Waals contribution ranges between 27 and 60 mJ/m2 and increases toward formation of a densely packed alkyl monolayer at C8, decreases due to the formation of a partial

MECHANICAL PROPERTIES OF CLAY MINERALS

121

alkyl bilayer with a minimum at C12, and increases toward formation of a densely packed alkyl bilayer at C20. The maximum cleavage energy of 126 mJ/m2 between C4 and C6 arises from a superposition of Coulomb and van der Waals contributions (Figure 4.15b). The minimum of the cleavage energy at C12 is related to a low Coulomb contribution and minimal van der Waals interactions in the partial alkyl bilayer. The cleavage energy reaches a local maximum for the densely packed bilayer at C20, and formation of a partial alkyl trilayer in the interlayer space causes a decrease at C22. We observe high lateral mobility of the head groups for short chains upon cleavage that decreases with increasing chain length. The exploitation of trends toward minimum cleavage energy DGM for low exfoliation free energies (Eq. (4.6)) is best possible for flat alkyl monolayers and flat alkyl bilayers in the montmorillonite gallery. When dialkylammonium surfactants, trialkylammonium surfactants, and other cationic surfactants with a higher load of organic material per head group are employed, interlayer structures increasingly part from a flat-on backbone conformation toward a tilted conformation [29,32] and the minimization of cleavage energies below 40 mJ/m2 becomes increasingly difficult. 4.3.3 Surface Energy The cleavage energy and the surface energy differ due to reference to different processes. The cleavage energy refers to the cleavage process and initial interlayer packing, the possible separation of Coulomb interactions between two layers upon cleavage, as well as surface reconstruction during cleavage contribute to its magnitude. The surface energy refers to surfaces that are already cleaved and measurements of contact angles do not take into account details of the separation. Therefore, the cleavage energy exhibits large fluctuations as a function of surfactant chain length and interlayer structure while the surface energy and the surface free energy (surface tension) are relatively insensitive to the interlayer structure and depend weakly on the chain length. The cleavage energy and the surface tension (¼surface energy minus 2–5 mJ/m2 for entropic contributions), as far as experimental data are available, are illustrated for montmorillonite of low CEC in Table 4.1. In the absence of strong Coulomb interactions across the two layers, the cleavage energy and the surface tension converge for densely packed alkyl layers, as the environment of the alkyl chains before cleavage and after cleavage in contact with solvents for contact angle measurements is similar. In all other cases, the differences can be large.

4.4 MECHANICAL PROPERTIES OF CLAY MINERALS Understanding the mechanical behavior of composites at the nanoscale also depends significantly on agreement on the tensile, shear, and bending moduli of aluminosilicate layers [37]. There has been considerable uncertainty across orders of magnitude [41], however, simulations using DFT and classical force field models have recently shown that nanoscale elastic constants are in good agreement with

122

g (expt) eS (sim)d Std. dev.

60 133e 5

None

36.2 2.5

C2

24.6 3.0

C4 46.4 35.8 2.0

C6 44.9 46.2 2.0

C8 44.9; 43.6 45.3 2.0

C10 c

Chain Length

42.3 41.2 2.0

C12 42.4 32.9 2.1

C14 39.1 2.3

C16

42.2 5.0

C18

51.4 3.0

C20

52.4 3.0

C22

In the absence of residual Coulomb contributions, the cleavage energy exhibits fluctuation as a function of the interlayer environment and approximates steady values of the surface tension in the case of densely packed alkyl interlayers. Reprinted from Ref. [34] with permission by the American Chemical Society. a Na-montmorillonite without surfactant modification for comparison. b CEC 68 mequiv./100 g. Data from Ref. [38]. c Ref. [22]. d CEC 91 mequiv./100 g. Data from Figure 4.14 Note that cleavage free energies are 2–5 mJ/m2 lower than cleavage energies (see text). e Ref. [27] and experiment.

b

a

TABLE 4.1 Experimentally Determined Surface Tension g and Computed Cleavage Energies « S (mJ/m2) for H3N þCnH2nþ1-Modified Montmorillonite

REFERENCES

123

TABLE 4.2 Nanomechanical Data (in GPa) for Layered Clay Minerals without Surface Modification

Ex Ey Ez K Gxy Gxz Gyz

Pyrophyllite (CEC 0)

Montmorillonite (CEC 91)

Montmorillonite (CEC 143)

Mica (CEC 251)

160 [1] 160 [4] 38 (1) 37 (1) 71 [1] 3–6 [0.2]

160 [3] 160 [4] 32 (1) 29 (1) 71 [1] 2–2.5 [0.25]

160 [3] 160 [20] 60 (1) 43 (1) 71 [3] 3.6–5 [0.25]

160 [6] 160 [20] 59 (1) 59 (2) 71 [3] 17 [1] 15 [1]

Young’s moduli E, bulk moduli K, and shear moduli G as a function of CEC (in mequiv./100 g) at low stress. The failure strength is indicated in square brackets where applicable, and the reference stress is indicated in round brackets for stress-dependent quantities. The reliability is approximately 5%. (Reprinted in simplified form from Ref. [37] with permission by the American Chemical Society.)

measured macroscale values. An overview of the most important elastic and shear constants is given in Table 4.2. The in-plane Youngs modulus Ex and Ey of 160 GPa as well as the in-plane shear modulus Gxy are essentially independent of CEC and of surface modification. Surface modification, however, affects the perpendicular Young’s modulus Ez, the bulk modulus K that is mostly influenced by Ez, as well as the perpendicular shear moduli Gxz and Gyz with the associated shear strengths. Unmodified, solvent-free minerals are of shear strength of 0.2–1 GPa parallel to the layers and a larger lateral stress leads to lateral shear flow between the layers. Upon organic surface modification, the perpendicular Young’s modulus and lateral shear moduli become similar to those of surfactants or short polyolefins, that is, on the order of 0.1 GPa or less [42]. Precise mechanical data for organically modified layered silicates are, to our knowledge, not yet available.

REFERENCES 1. Lagaly, G.; Weiss, A. Arrangement and orientation of cationic tensides on silicate surfaces. 2. Paraffin-like structures in alkylammonium layer silicates with a high layer charge (mica). Koll. Z. Z. Polym., 237, 364–368 (1970). 2. Lagaly, G.; Weiss, A. Arrangement and orientation of cationic tensides on silicate surfaces. 4. Arrangement of alkylammonium ions in low-charged silicates in films. Koll. Z. Z. Polym., 243, 48–55 (1971). 3. Lagaly, G. Kink-block and gauche-block structures of bimolecular films. Angew. Chem. Int. Ed., 15, 575–586 (1976). 4. Vaia, R. A.; Teukolsky, R. K.; Giannelis, E. P. Interlayer structure and molecular environment of alkylammonium layered silicates. Chem. Mater., 6, 1017–1022 (1994). 5. Vaia, R. A.; Giannelis, E. P. Lattice model of polymer melt intercalation in organically modified layered silicates. Macromolecules, 30, 7990–7999 (1997).

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6. Vaia, R. A.; Giannelis, E. P. Polymer melt intercalation in organically modified layered silicates: model predictions and experiment. Macromolecules, 30, 8000–8009 (1997). 7. Hackett, E.; Manias, E.; Giannelis, E. P. Molecular dynamics simulations of organically modified layered silicates. J. Chem. Phys., 108, 7410–7415 (1998). 8. Hayes, W. A.; Schwartz, D. K. Two-stage growth of octadecyltrimethyl-ammonium bromide monolayers at mica from aqueous solution below the Krafft point. Langmuir, 14, 5913–5917 (1998). 9. Brovelli, D.; Caseri, W. R.; Hahner, G. Self-assembled monolayers of alkylammonium ions on mica: direct determination of the orientation of the alkyl chains. J. Coll. Interface Sci. 216, 418–423 (1999). 10. Osman, M. A.; Seyfang, G.; Suter, U. W. Two-dimensional melting of alkane monolayers ionically bonded to mica. J. Phys. Chem. B, 104, 4433–4439 (2000). 11. Pospisil, M.; Capkova, P.; Merinska, D.; Malac, Z.; Simonik, J. Structure analysis of montmorillonite intercalated with cetylpyridinium and cetyltrimethylammonium: molecular simulations and XRD analysis. J. Coll. Interface Sci.; 236, 127–131 (2001). 12. Osman, M. A.; Ernst, M.; Meier, B. H.; Suter, U. W. Structure and molecular dynamics of alkane monolayers self-assembled on mica platelets. J. Phys. Chem. B, 106, 653–662 (2002). 13. Kuppa, V.; Manias, E. Computer simulation of PEO/layered silicate nanocomposites: 2. Lithium dynamics in PEO/Li þ montmorillonite intercalates. Chem. Mater., 14, 2171–2175 (2002). 14. Heinz, H.; Castelijns, H. J.; Suter, U. W. Structure and phase transitions of alkyl chains on mica. J. Am. Chem. Soc., 125, 9500–9510 (2003). 15. Pospisil, M.; Kalendova, A.; Capkova, P.; Simonik, J.; Valaskova, M. Structure analysis of intercalated layer silicates: combination of molecular simulations and experiment. J. Coll. Interface Sci., 277, 154–161 (2004). 16. Heinz, H.; Suter, U. W. Surface structure of organoclays. Angew. Chem. Int. Ed., 43, 2239–2243 (2004). 17. Osman, M. A.; Ploetze, M.; Skrabal, P. J. Structure and properties of alkylammonium monolayers self-assembled on montmorillonite platelets. Phys. Chem. B, 108, 2580–2588 (2004). 18. Heinz, H.; Paul, W.; Binder, K.; Suter, U. W. Analysis of the phase transitions in alkyl-mica by density and pressure profiles. J. Chem. Phys., 120, 3847–3854 (2004). 19. Zeng, Q. H.; Yu, A. B.; Lu, G. Q.; Standish, R. K. Molecular dynamics simulation of the structural and dynamic properties of dioctadecyldimethyl ammoniums in organoclays. J. Phys. Chem. B, 108, 10025–10033 (2004). 20. Zhu, J. X.; He, H. P.; Zhu, L. Z.; Wen, X. Y.; Deng, F. J. Characterization of organic phases in the interlayer of montmorillonite using FTIR and 13C NMR. J. Coll. Interface. Sci., 286, 239–244 (2005). 21. Lewin, M.; Mey-Marom, A.; Frank, R. Surface free energies of polymeric materials, additives, and minerals. Polym. Adv. Technol., 16, 429–441 (2005). 22. Osman, M. A., Rupp, J. E. P.; Suter, U. W. Gas permeation properties of polyethylenelayered silicate nanocomposites. J. Mater. Chem., 15, 1298–1304 (2005). 23. Pandey, R. B.; Anderson, K. L.; Heinz, H.; Farmer, B. L. Conformation and dynamics of a self-avoiding sheet: bond-fluctuation computer simulation. J. Polym. Sci. B, 43, 1041–1046 (2005)

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24. Heinz, H.; Koerner, H.; Anderson, K. L.; Vaia, R. A.; Farmer, B. L. Force field for micatype silicates and dynamics of octadecylammonium chains grafted to montmorillonite. Chem. Mater., 17, 5658–5669 (2005). 25. Greenwell, H. C.; Harvey, M. J.; Boulet, P.; Bowden, A. A.; Coveney, P. V.; Whiting, A. Interlayer structure and bonding in nonswelling primary amine intercalated clays. Macromolecules, 38, 6189–6200 (2005). 26. He, H. P.; Galy, J.; Gerard, J. F. Molecular simulation of the interlayer structure and the mobility of alkyl chains in HDTMAþ /montmorillonite hybrids. J. Phys. Chem. B, 109, 13301–13306 (2005). 27. Heinz, H.; Vaia, R. A.; Farmer, B. L. Interaction energy and surface reconstruction between sheets of layered silicates. J. Chem. Phys., 124, 224713 (2006). 28. Jacobs, J. D.; Koerner, H.; Heinz, H.; Farmer, B. L.; Mirau, P.; Garrett, P. H.; Vaia, R. A. Dynamics of alkyl ammonium intercalants within organically modified montmorillonite: dielectric relaxation and ionic conductivity. J. Phys. Chem. B, 110, 20143–20157 (2006). 29. Heinz, H.; Vaia, R. A.; Krishnamoorti, R.; Farmer, B. L. Self-assembly of alkylammonium chains on montmorillonite: effect of chain length, headgroup structure, and cation exchange capacity. Chem. Mater., 19, 59–68 (2007). 30. Fermeglia, M.; Pricl, S. Multiscale modeling for polymer systems of industrial interest. Prog. Org. Coat., 58, 187–199 (2007). 31. Scocchi, G.; Posocco, P.; Fermeglia, M.; Pricl, S. Polymer–clay nanocomposites: a multiscale molecular modeling approach. J. Phys. Chem. B, 111, 2143–2151 (2007). 32. Heinz,H., Vaia, R. A.; Farmer, B. L. Relation between packing density andthermal transitions of alkyl chains on layered silicate and metal surfaces. Langmuir, 24, 3727–3733 (2008). 33. Heinz, H.; Vaia, R. A.; Koerner, H.; Farmer, B. L. Photoisomerization of azobenzene grafted to montmorillonite: simulation and experimental challenges. Chem. Mater., 20, 6444–6456 (2008). 34. Fu, Y. T.; Heinz, H. Cleavage energy of alkylammonium-modified montmorillonite and the relation to exfoliation in nanocomposites: influence of cation density, head group structure, and chain length. Chem. Mater., 22, 1595–1605 (2010). 35. Fu, Y. T.; Heinz, H. Structure and cleavage energy of surfactant-modified clay minerals: Influence of CEC, head group, and chain length. Phil. Mag., 90, 2415–2424 (2010). 36. Heinz, H.; Suter, U. W. Atomic charges for classical simulations of polar systems. J. Phys. Chem. B, 108, 18341–18352 (2004). 37. Zartman, G. D.; Liu, H.; Akdim, B.; Pachter, R.; Heinz, H. Nanoscale tensile, shear, and failure properties of layered silicates as a function of cation density and stress. J. Phys. Chem. C, 114, 1763–1772 (2010). 38. Giese, R. F.; van Oss, C. J. Colloid and Surface Properties of Clays and Related Minerals, Dekker, New York, 2002. 39. Yariv, S.; Cross, H (eds.). Organo-Clay Complexes and Interactions, Dekker, New York, 2002. 40. Kamal, M. R.; Calderon, J. U.; Lennox, R. B. Surface energy of modified nanoclays and its effect on polymer/clay nanocomposites. J. Adhesion Sci. Technol., 23, 663–688 (2009). 41. Chen, B.; Evans, J. R. G. Elastic moduli of clay platelets. Scripta Mater., 54, 1581–1585 (2006). 42. Brandrup, J.; Immergut, E. H.; Grulke, E. A. (eds.). Polymer Handbook, Wiley, New York, 1999.

CHAPTER 5

CHEMISTRY OF RUBBER–ORGANOCLAY NANOCOMPOSITES LUCA GIANNINI ATTILIO CITTERIO MAURIZIO GALIMBERTI DAFNE COZZI

5.1 INTRODUCTION The use of organic cations, in particular of quaternary ammonium salts, to improve the disperdibility of mineral fillers such as talc, kaolin, and bentonite in low polarity plastics and rubbers dates back to the 1980s [1], while the effect of these organic cations on the sulfur vulcanization of rubber compounds was investigated even before that time; quaternary ammonium and phosphonium salts were recognized as secondary accelerators in the vulcanization of unsaturated elastomers with a sulfur system [2], with the exception of aryl-substituted ammonium, which were found to have low accelerating effect, or even a retarder effect, which might be interpreted reminding that aryl-substituted amines are a well-known family of radical scavengers, commonly used to stabilize unsaturated rubber toward oxidation. The vulcanization of rubber nanocomposites comprising organoclays, that is, swelling clays modified with organic cations, is discussed in Chapter 7; as organoclays are compatibilized with low polarity polymers using alkyl substituted ammonium, phosphonium, and imidazolinium salts, they are in general expected to have an accelerating effect in sulfur-based vulcanization. As a matter of fact, most authors observed that organoclays accelerated the rate of vulcanization, and interpreted this behavior as an effect of the release of basic species (amines) from the organoclays. In most of the papers reviewed in Chapter 7, the organoclays included ammonium salts and the corresponding alkyl-substituted amines are well known as secondary accelerators. Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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128

CHEMISTRY OF RUBBER–ORGANOCLAY NANOCOMPOSITES

In the case of organoclays containing ammonium salts obtained by protonation of an amine (e.g., a primary amine such as octadecylamine), the release of the corresponding amine in vulcanization conditions appeared likely. In the case of organoclays featuring fully substituted organic cations, that is, organic cations that cannot be converted to neutral species by deprotonation of the heteroatom (e.g., quaternary ammonium, phosphonium or fully substituted imidazolium cations), a more complex situation is envisageable. The observed reactivity might be due to the generation of basic species by thermal decomposition processes, but might also be due to the intrinsic reactivity of the organic cations with the vulcanization active species and the counter anions associated to these species. This chapter discusses the chemistry of organoclays in the frame of a sulfur vulcanized rubber compound, both in terms of thermal decomposition of the organic cations, and in terms of their intrinsic reactivity. 5.2 ORGANIC CATION DECOMPOSITION IN SALTS, ORGANOCLAYS AND POLYMER NANOCOMPOSITES The thermal decomposition of polymer nanocomposites comprising organoclays modified with fully substituted organic cations, that is, organic cations that cannot be converted to neutral species by deprotonation of the heteroatom (e.g., quaternary ammonium, phosphonium, or imidazolium) was investigated by various authors [3–34], considering a number of aspects: (i) the relation to the thermal decomposition of the organoclay and the organic salt whose cationic part was used for the modification of the clay; (ii) the effect of the degree of organic cation substitution and the additional presence of organic salts; (iii) the presence of nucleophiles in the polymer nanocomposite. Table 5.1 summarizes references to papers dealing with the thermal decomposition of organoclays, the precursor organic cations salts and in some cases polymer nanocomposites of such organoclays. 5.2.1 Experimental Techniques The experimental techniques most used to investigate the thermal decomposition of the organoclays were thermogravimetric ones, often coupled to IR, MS, and GC for the identification of decomposition volatile products; thermal analysis, XRD diffraction spectra, and TEM observation completed in some cases the characterization. One of the most relevant parameters in thermal analysis to describe the stability of organoclays and related species appears to be the “onset decomposition temperature,” that is, the temperature at which a significant mass loss (not due to evaporation of adsorbed water) is first observed. It should be noted that, as standardized methods for the thermal analysis of organoclays do not exist, different experimental protocols were adopted, and as a result the absolute data reported in different papers are often not well comparable. Despite this fact, general trends are consistent and allow to draw solid conclusions.

129

P-C16,3Ph; N-VBz; N-C16, VBz; N-C16, PhCH2OH

N-C; N-T; N-C12; N-C18; N-2C; N-2HT

N-C; N-T; N-C12; N-C18

Organic Cation Type




As Salt

Onset decomposition temperature [N2,10 C/min] increased by >50 C for phosphonium modified organoclay versus ammonium. Degradation products: alkenes, amines, aldehydes, PPh3. Mechanistic hyp.: Hoffmann elimination. Polystyrene nanocomposites stabilized in all cases




Onset decomposition temperature [N2, 20 C/min] organoclays: 218–222 C; degradation products: alkanes, alkenes, amines, haloalkanes, aldehydes

Key Conclusions

Onset decomposition temperature [N2, 2 C/min] organoclays: 155–180 C ¼ 15–25 C lower versus salts. Washing/ extracting organoclay: d-spacing decreases and thermal stability increases. d-Spacing increases up to 280 C, disorder increases above 200 C. Degradation products: alkanes, alkenes, amines, haloalkanes, aldehydes. Mechanistic hyp.: Hoffmann elimination (dominant in organoclay) and SN2 (dominant in salt). Organoclay has complex decomposition with 4 peaks in DTG, salt only 1 peak Polystyrene

In Organoclay in Polymer







In Organoclay

Thermal Decomposition Studied

TGA/FTIR

TGA/MS, High Resolution DTG, GC/MS, TG/FTIR, XRD

TG/MS, SDT TGA/FTIR

Analytical Techniques

(continued )

5

4

3

Reference

TABLE 5.1 Summary of Papers Dealing with Organic Cation Decomposition in Salts, Organoclays and Polymer Nanocomposites

130

In Organoclay







As Salt







IMZ-Me2,C3; IMZ-Me2,C6; IMZ-Me2,C10; IMZ-Me2,C16; N-2C18; N-2HT

P-3PhC12; P-3C4C14; P-3C4C16; P-3C4C18; P-4Ph; P-4C8; N-4C8

In Organoclay in Polymer

Thermal Decomposition Studied

Organic Cation Type

TABLE 5.1 (Continued )

Onset decomposition temperature [N2, 2 C/min] phosphonium organoclays: 190–230 C ¼ 70–80 C lower versus phosphonium Br salts. Organoclays have complex decomposition with 4 peaks in DTG, salts only 1 peak. d-Spacing increases wth T up to 370 C (melting then pressure build-up), then collapses. Degradation products: alkenes, haloalkanes (salts), phosphines, phosphine oxide (organoclays). Mechanistic hyp.: b-elimination, nucleophilic substitution at C (salts) and P (organoclay)

Onset decomposition temperature [N2, 10 C/min] increased by >50 C for imidazolium modified organoclay versus quat. ammonium. Residual chloride anion decreased stability

Key Conclusions

TGA, DTG, Pirolysis/G C-MS (XRD, FTIR) versus T

TGA-DTA, TDMS

Analytical Techniques

7

6

Reference

131

N-C18 P-Bu4 P-Bu3Ph P-C16-Bu3

Onset decomposition temperature [N2, 10 C/min] phosphonium salts: 255–287 C; ammonium salt: 202 C; organoclay phosphonium: 225–302 C; organoclay ammonium: 158–222 C as a function of ammonium cation concentration: 0.8 CEC most stable (vs. 1 and 1.2 CEC)

Washing organoclay: some 2HT is removed, d-spacing is reduced, thermal stability increases, polarity decreases. Swelling in solvents does not involve d-spacing change




N-2HT




Onset decomposition temperature [air, 5 C/min] Organoclay: 213–253 C as a function of halide contamination: most stable when halide free. Degradation products: alkenes, amines. Mechanistic hyp.: Hoffmann elimination main path




N-2C18




Onset decomposition temperature [N2, 10 C/min] imidazolium salts: 225–400 C; >100 C effect of anion (Cl, Br low stability); linear short substituents most stable. Organoclays more stable than halide salts: onset dec. T ¼ 280–340 C Mechanistic hyp.: SN2, SN1 (branched substituents)







IMZ-Me2,C3; IMZ-Me2,C4; IMZ-Me,C4; IMZ-Me,C10; IMZMe,C16; IMZ-Me2, C16; IMZ-Me2, eicosyl; IMZ-Me2, ethylbenzene; IMZ-Me,ethyl; IMZ-Me2,C3; IMZ-Me2,isobutyl; IMZ-Me2,C10; IMZ-Me2,allyl

High resolution TGA, DTG, DSC, XRD

TGA, DTG, XRD, SWELLING

TGA MS

TGA, TGA FTIR MS

(continued )

19

18

12

11

132




As Salt

Isoprene rubber







In Organoclay in Polymer







In Organoclay

Thermal Decomposition Studied

Review - Introduction of layered silicates into polymer matrix causes an increase in thermal stability, function of clay dispersion, attributed to reduction in gas permeability and heat conduction, easier char formation, possibly due to catalytic effects

Onset decomposition temperature [N2, 5–20 C/ min] rank: masterbatches in IR< (5–7 C); organoclays < (15–30 C) salts Benzyl group least stable. Degradation products: alkenes, haloalkanes (for salt þ organoclay if impure), amines, þ (for IR MB containing stearic acid) long chain carboxylic esters. Mechanistic hyp.: SN2 dominating mechanism, O on periphery of clay might be alkylated

Onset decomposition temperature [N2, 10 C/min]: 160–198 C; benzyl group most labile, followed by (CH2)2OH. Degradation products: alkanes, alkenes, aldehydes, acids, amines. Mechanistic hyp.: Hoffmann, SN2, thermal degradation of tallow

Key Conclusions

TEM TGA DTG GC-MS

TGA, FTIR

Analytical Techniques

33

31

21

Reference

Bz ¼ benzyl; C ¼ coco; T ¼ tallow; HT ¼ hydrogenated tallow; N-R ¼ Me3N-R þ ¼ quaternary ammonium cation with “R” group and three methyl groups; P-R ¼ Me3P-R þ quaternary phosphonium cation with R group and three methyl groups; IMZ-Me2,C16 ¼ imidazolinium substituted with a C16 alkyl chain and two methyl groups

Ammonium, imidazolium, phosphonium, oligomeric cations

N-2T N-T,Bz N-T.C8

N-HT,Bz; N-2HT; N-HT,C8; N-T, (CH2)2OH

Organic Cation Type

TABLE 5.1 (Continued )

ORGANIC CATION DECOMPOSITION IN SALTS

133

5.2.2 Decomposition of Organoclays Versus Precursor Organic Cation Salts The most striking difference between the thermal behavior of organoclays with respect to their precursor organic cation salts is that the latter show a simple behavior, represented by a single peak in DTG, while organoclays show a complex behavior, represented by at least four peaks, spanning an interval of approx. 600 C, as reported in Figures 5.1 and 5.2 [4]. Figure 5.2 also illustrates the different results obtained using different TGA techniques. The “onset decomposition temperature” of organoclays is typically close to that of their precursor organic salts, but the decomposition extends at, and beyond, the temperature range where dehydration of the aluminosilicate occurs (450–700 C). It is evident that the presence of the inorganic clay induces different reaction paths of the organic cations with respect to the precursor salts, and in particular that the initial products undergo successive secondary reactions. The presence of oxygen and metal species in the montmorillonite structure may serve as catalysts to enable the oxidative cleavage of alkenes to produce aldehydes, acids and CO2 at elevated temperature. The weight loss observed in organoclays between 700 and 1000 C is attributed to the decomposition of a carbonaceous residue. A number of papers compare the thermal stability of the organic cations with their respective organoclays, reporting both a decrease and an increase of the “onset decomposition temperature”; these data can be interpreted on the basis of the decomposition mechanisms of organic cations as discussed in detail in this chapter, the dominant mechanisms in the decomposition of organic cations involve their interaction with a Lewis base acting as a nucleophile or abstracting a proton. When the pristine salt contains a rather active anion (e.g., chloride), the thermal stability of 2.5 2.0 Deriv. weight (% / ºC)

2048 × 10

1.5 TMO

1.0 0.5 0.4

–0.0 0

100

200 300 Temperature (ºC)

400

500

FIGURE 5.1 DTG curves of trimethyloctadecylammonium chloride (TMO) and the corresponding montmorillonite organoclay. Reprinted with permission from Ref. [4].

134

CHEMISTRY OF RUBBER–ORGANOCLAY NANOCOMPOSITES

0.25

(a)

Deriv. weight (% / ºC)

0.15

0.05

–0.00

–0.15

–0.25

0

200

400

600

800

1000

Temperature (ºC) 0.15 (b)

Deriv. weight (% / ºC)

0.10 0.05 0.00

–0.05

–0.10

–0.15 0

200

400

600

800

1000

Temperature (ºC)

FIGURE 5.2 DTG curves from conventional and high-resolution TGA of pristine montmorillonite (a) and a montmorillonite organoclay exchanged with trimethylcocoammonium chloride (b). Reprinted with permission from Ref. [4].

the organic cation (as measured by the “onset decomposition temperature”) is decreased in the organoclay if a significant amount of such anion is present into the organoclay [4,7,19,31], while is increased if the reactive anion is not present in the organoclay [19]. Removal of excess onium halide salts from organoclays was reported to increase the thermal stability of organoclays [4,12,18]. A major effect of the anion nucleophilicity was demonstrated on the thermal stability of

MECHANISM OF THERMAL DECOMPOSITION OF ORGANOCLAYS

135

alkyl-imidazolium salts [11]; the thermal stability of the corresponding alkylimidazolium organoclays (in the absence of reactive anions) was found to be comparable to that of the salts with the least reactive anions. The advantage of using phosphonium instead of ammonium salts is that they give rise to an increase in the degradation temperature of the organoclay from 200–250 to 300–350 C [35]. 5.3 MECHANISM OF THERMAL DECOMPOSITION OF ORGANOCLAYS Mechanistic hypothesis on the thermal decomposition of organoclays are based on the reaction products that can be detected as a consequence of the decomposition process. Thermal decomposition of alkyl quaternary ammonium cations is known to proceed via Hofmann elimination or SN2 nucleophilic substitution reaction (Figure 5.3) at a temperature as low as 155 C, irrespective of the molecular structure [4]. Generally, the first decomposition mode is preferred with hard anions, whereas with more polarizable anions the latter process prevails. The main decomposition product expected from the thermal degradation of organoclays modified with quaternary ammonium salts is thus represented by tertiary amines; this is confirmed by analysis of by-products. It was reported that, especially at lower temperatures, amines generated by abstraction of the residue with the lowest N–C bond strength are prevalent; benzyl residue is preferentially lost, than long-chain alkyl, while methyl group is the most stable [31]. On the other hand, it was also reported that in a Nylon nanocomposite, methyl group was preferentially lost from dimethylditalloyl ammonium salt [36]. In the case of organoclays substantially free of halides contamination, 1-olefins corresponding to alkyl substituents on heteroatom were formed together with tertiary amines as the prevalent decomposition products of quaternary ammonium [3–5,12] and phosphonium [5,7] salts, pointing to the prevalence of Hofmann elimination reaction over SN2 (and any other mechanisms envisageable for phosphonium salts) [7]. At high decomposition temperatures, oxidized species and smaller fragments, resulting from secondary reactions, appear [4,5].

R3

H + N C R1 C H R R2 H2 X– X– R3 H + R1 N C C H R R2 H2

X

C H2

H2 C

N R3 R + R1 R2

H N HX + H C C R + R1 R3 2 R2

FIGURE 5.3 Formation of tertiary amines by SN2 substitution (upper scheme) or by Hofmann elimination (lower scheme).

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CHEMISTRY OF RUBBER–ORGANOCLAY NANOCOMPOSITES

Haloalkanes (RX, X ¼ Cl, Br) are always found when organoclays contain a significant amount of halide ions, typically coming from an excess of organic halide salt [3,4,7,31]. The mechanism of reaction is supposed to be a SN2 attack of the halide ion (X) to the a-carbon of organic residues (R) bound to heteroatom (Y) in organic cations, that is, the reverse of the quaternary onium synthesis, yielding RX and RxY. In support of this hypothesis is also the finding that no 2-chloroalkanes (but some 1-chloroalkanes) were revealed by GC in the decomposition products of organoclays modified by a slight excess of quaternary ammonium chlorides [31]. Thermal decomposition of masterbatches comprising quaternary ammonium organoclays, synthetic isoprene rubber, and stearic acid was studied [31] in conditions close to that of rubber vulcanization, heating the masterbatch samples at 170 C for 1 h while collecting volatiles, which were then analyzed by GC–MS (Table 5.2). Degradation of masterbatches with organoclays in IR rubber mainly yielded long-chain tertiary amines, long-chain methyl esters, 1-chloroalkanes and long-chain 1-olefins. Tertiary dimethylalkylamines are as expected the prevalent products. Less obvious is the formation of long-chain carboxylic esters. However, the presence of benzyl esters in masterbatches 4 and 5 suggested that carboxylic acid present in the masterbatches were alkylated by quaternary ammonium cations or by 1-chloroalkanes. Surprisingly, the amount of long-chain 1-olefins was always very low, despite the fact that Hofmann elimination is suggested as the major decomposition process of quaternary ammonium cations; this fact might be explained (i) by the low temperature at which the experiment was conducted, corresponding to the “first event” in DTG curve of organoclays featuring quaternary ammonium salts, usually involving only the decomposition of excess ammonium halide [4] (with predominant formation TABLE 5.2 Representative Organic Species Detected by GC-MS in the Thermal Decomposition of IR/Organoclay Masterbatches (wt %) (Adapted from Ref. [31]) Masterbatches in IR þ Stearic Acid treta Ammonium cation Max mixing T ( C) Products Methyloctadecanoate N,N-Dimethyloctadecylamine 1-Chlorooctadecane Methylhexadecanoate N,N-Dimethyltetradecylamine 1-Chlorohexadecane 1-Octadecene Benzyl hexadecanoate Benzyl octadecanoate N,N-Dimethyl-2-ethylhexylamine a

Retention time of identified products.

23.25 23.18 23.00 21.60 21.40 21.14 22.56 27.16 29.35 17.57

2

3

4

5

6

7

2T 130

2T 174

T, Bz 160

T, Bz 195

T, C8 165

T, C8 192

25% 36% 11% 3% 16% 9% 0.4%

21% 36% 14% 2% 17% 10% 0.2%

1% 57% 1% 0.3% 38% 1% 0.1% 1% 1%

2% 62% 1%

49% 10% 3% 7% 6% 8% 6%

48% 18% 5% 7% 8% 6% 2%

11%

5%

32% 1% 0.1% 0.4% 0.5%

ROLE OF ORGANIC CATIONS IN ORGANOCLAYS AS RUBBER VULCANIZATION ACTIVATORS

137

R3 R1

N+ R2

Si O Si O Si

R

R1

N R2

R3

+ R

O-

Si O Si O Si O

FIGURE 5.4 SN2 substitution of quaternary ammonium cation by oxygen anion of clay surface. Reprinted with permission from Ref. [4].

of haloalkanes); (ii) by the presence in the masterbatch of carboxylic acids (ubiquitous in rubber technology), which seem to act as nucleophiles, in concurrence with halide anions. The presence of tertiary amines containing talloyl and 2-ethylhexyl radicals but not the benzyl radical indicates that the type of substituent of the nitrogen atom deeply affects the fragmentation reactions of the ammonium salt. Thermal decomposition of organic ammonium cations predominantly occurs at the site with the smallest N–C bond dissociation energy. In this respect, it can be easily interpreted the very minor presence of tertiary amines derived from loss of a methyl groups, that are indeed strongly bonded to the nitrogen atom. Moreover, because the sum of alkene and 1-chloroalkanes was in defect with respect to tertiary amines other nucleophilic agents, either of low molecular mass or present on the surface of clay, might be involved in SN2 substitution reactions. Similarly, the formation of long-chain carboxylic acid methyl esters could be explained by the nucleophilic substitution to the alkylammonium cation by the corresponding carboxylate anion (fatty acids were in fact used in all masterbatches). These considerations suggest the possibility that nucleophilic oxygen atoms in the periphery of the clay gallery could undergo the substitution with a net increase of hydrophobicity of the inorganic material (Figure 5.4). The decomposition of imidazolium cations was reported to proceed mainly by SN2, except for branched alkyl substituents where SN1 mechanism was suggested [11]. All these data suggest that thermal degradation of the modifier will not only alter the carefully tailored surface compatibility but also the resulting products may play a major and yet to be determined role in the formation of exfoliated nanostructures and the physical characteristics of the final nanocomposite. 5.4 ROLE OF ORGANIC CATIONS IN ORGANOCLAYS AS RUBBER VULCANIZATION ACTIVATORS The reaction mechanism of the rubber vulcanization process is still under debate due to the complex interplay of several elemental steps involved. In particular, relevant is the role of accelerators/activators interaction, each component influencing the reactivity of the other. Proposed mechanisms have ranged from radical to ionic,

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CHEMISTRY OF RUBBER–ORGANOCLAY NANOCOMPOSITES

Accelerators + activators Active accelerator complex Sulfur donor + activators Sn Active sulfurating agent R-H Rubber-bound intermediate (RSyX) Initial polysulfide cross-links (RSxR) Cross-link shortening/additional crosslinking Cross-link destruction/chain modification S-S bond interchange

Final vulcanizate network

FIGURE 5.5 General reaction scheme for the sulfur vulcanization of rubbers (R represents the organic chain and X is the accelerator residue).

and several researchers have concluded that both mechanisms are operative. There is now a widespread agreement about the basic steps in accelerated sulfur vulcanization, proposed by Morrison and Porter in 1984 [37] (Figure 5.5). The cationic part of organoclays can play, in this picture, a specific role not only as thermal generator of redox active neutral species (i.e., tertiary amines) but also as surface active ingredients participating in the interphase transfer reactions developing in the course of vulcanization. In particular, it can work as activator increasing the mobility of sulfur (and other heteroatoms) accelerating anionic species. There are indications that catalysts of interphase transfer, like combinations of quaternary ammonium salts and sodium hydroxide, take part in the degradation reactions of polysulfide units in the regeneration of sulfur vulcanizates [38] and that QAS sulfides and polysulfides are easily involved in nucleophilic substitution to S–S and S–N containing molecules. The so-called multifunctional additives (MFA, whose general structure is [R0 NH2(CH2)3NH3]2 þ 2[C17H33COO]  ) were first reported to function in their own right as simultaneous activator and accelerator for sulfur vulcanization [39]. These cationic surfactants imparted to a rubber compound a combination of improved properties in a MFA/sulfur-NR/SBR tire sidewall type rubber formulation acting as a cure activator and accelerator. Their use eliminated the need of a mould release agent, decreased viscosity and improved flow, promoted wetting and dispersion of fillers, enhanced tack and auto adhesion, and performed as a lubricant, improving extrusion, calendaring and molding. It was reported that it was not necessary to have ZnO and stearic acid present in the formulation [40]; however it was found that the inclusion of ZnO in a MFA containing recipe increased the scorch time drastically and eliminated cure reversion [41], as expected upon introduction of “normal” ZnO levels. The mechanistic details of all these activation are still unclear, in particular the initiation step. However, the proximity of the Lewis base sites and the basic aluminosilicate surface to the intercalated alkyl quaternary ammonium molecule, as previously indicated, is conducive to enhancing (lowering the energy of) the

139

ROLE OF ORGANIC CATIONS IN ORGANOCLAYS AS RUBBER VULCANIZATION ACTIVATORS

OH

HO

A–

O– H+

X–

HO

OH

OH

A-H

X–

O– H+

O– H+

X-H A-H

A–

R-A

2C13

X-H O– HO

R-A

H+

H+ A-H

O–

X–

OH

X–

A-H OH

OH

H+ O–

A–

OH

FIGURE 5.6 Schematic representation of possible structures of C18-QATMs with adsorbed anion and neutral molecules on montmorillonite clay.

Hofmann elimination reaction but also to exchanging protons with adsorbed acidic species. It is well documented that the interlayer of organically modified clay can host several low and medium molecular weight molecules owing to the free space left from the adsorbed organic cation. In fact, the degree of the surface coverage by alkylammonium cations in montmorillonite (M) (overall surface area of 750 m2/g and cation-exchange capacity is 90 mequiv./100 g) was evaluated 0.86, with a layer area per molecule available to QA of 0.68 nm2. A relevant volume is so open to the absorption of other neutral and charged species (Figure 5.6). The proton exchange with the basic aluminosilicate surface will improve the association of the cation with all anions, that is, the ones previously present and those arising from the deprotonation of acids. As shown by Fripiat and Stone [42], the water adsorbed on clay may be source of the mobile protons as dissociation degree of such adsorbed water is 106 higher than in bulk. The so formed organic ion pair can migrate into the nonpolar polymer bulk, increasing both the reactivity of anions toward SN2 substitutions and redox reactions and favoring, as phase transfer agent [43], the exchange of anions at the surface of solid additives present in masterbatches (i.e., ZnO). Figure 5.7 summarizes the acid–base equilibria and nucleophilic substitution reactions involved in the generation of a RX vulcanization initiator starting from the Q þ  OSiR3 organoclay and HA or HX acid molecules. The relative magnitude of the observed nucleophilic substitution and anioninduced redox initiation will depend on factors effecting mass transfer of the products, such as defect concentration, perfection of layer stacking, crystallite size, and packing density of the aggregate. Sulfur centered anions associated in the ion pair with the onium cation are the more plausible reactive species involved in the vulcanization process in the presence of organoclay. Both simple inorganic sulfide (S2  , HS  ) and polysulfides (Sx2  ) or . thiolate (RS  ) and organopolysulfides (RSx  ), but also radical anions (i.e., S  , . . S2 or S3 ), can play a specific role owing to their similar but not coincident

140

CHEMISTRY OF RUBBER–ORGANOCLAY NANOCOMPOSITES

Q+X– + RA

RX + Q+A–

Q+X– + RA

RX + Q+A–

(Polymeric phase) (Clay surface phase)

Q+OSi– + AH

HOSi– + Q+A–

Q+X– + HOSi–

HX + Q+OSi–

RX

Vulca initiation

FIGURE 5.7 Catalytic mechanism of rubber vulcanization promoted by onium clay (Q þ  OSi).

chemistry. Anionic species can be generated by deprotonation of thiols (RSH) and organosulfanes (RSxH) or by attack of strong nucleophiles on the S–S bond of organic sulfur cluster molecules present in the rubber masterbatch, that is, Nu  þ RSS-SSR ! RSS  þ RSS-Nu These last reactions are between the more important ones in the chemistry of polysulfur compounds and their rates depend on the strength of the nucleophile, the too weak ones R3N, OH  , S2O32  , and SCN  being rather unreactive [44]. Sulfur centered anions are strong nucleophiles and powerful reducing species, in turn easily involved in nucleophlic substitution and electron transfer processes, whereas polysulfide anions are less nucleophilic and show more complex redox chemistry [45]. Substitution reactions by sulfur anions on neutral sulfur cluster molecules (Sx, RSxR), ions (RSx  ) and sulfenamides (RS-NRR0 ) are fast also at moderate temperatures and occur by thiophilic attack on S atom because sulfur is known to be able to expand its octet into decet. The participation of p and d orbitals in these processes can lower the energy of activation, as for instance in the S8 dissolution in polysulfide solutions. S22  þ S8 ¼ S102  S102 þ S22  ¼ 2S62  ¼ S52  þ S72  Relatively few onium salts of polysulfide anions are known and isolated [46]. The best investigated are the heptasulfide salts (i.e., (Pr4N)2S7 [47] and (Ph4P)2S7 [48]). Moreover, polysulfide onium salts have been postulated as intermediates in PTC nucleophilic substitution reactions for the synthesis of organic polysulfanes [49], in sulfur removal from oil [50], and in polysulfide polymer preparation [51]. Polysulfide anions are also known to be thermally unstable [52] to homolytic dissociation (giving . . ubiquitous radical anion species, that is, the highly colored S2  or S3  ) and sensitive to molecular oxygen and other oxidants [53].

REFERENCES

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The reactions of sulfur centered anions with oxygen and other oxidant agents are fast and general processes owing to the low redox potentials of these species. In water sulfide ions have a reducing power higher than metal iron (Eh(S0/S2  ) ¼  0.48 V, NHE versus Eh(Fe2 þ /Fe) ¼ 0.44 V) whereas polysulfides are less reducing (depending on sulfur atom bonded) (Eh ¼  0.42–0.33) and thiolate and oganopersulfide anions are even weaker reducing agents. Mono electron transfer reactions from sulfur centered anions to oxidant agents are easier for polysulfides to give perthiyl radicals . . (RSx , x  2) than thiol to give thiyl radicals (RS ), being the first stabilized by partial double or p-bond formation, an effect absent in the latter, whereas the monoelectronic . oxidation of polysulfide anions mainly affords S3  radical anion [54]. All these paramagnetic species can play a role in the initiation, propagation, and termination steps of sulfur vulcanization. REFERENCES 1. Weber, K. E.; Mukamal, H. Rubber compositions comprising phillosilicate minerals, silanes and quaternary ammonium salts. U.S. Patent 4,431,755 (1984). 2. Scheele, W.; Petry, G. Phosphonium halides as vulcanization accelerators. U.S. Patent 3,514,430 (1970), and references therein. 3. Xie, W.; Gao, Z.; Liu, K.; Pan, W.; Vaia, R.; Hunter, D.; Singh, A. Thermal characterization of organically modified montmorillonite. Thermochim. Acta, 367–368 339–350 (2001). 4. Xie, W.; Gao, Z.; Pan, W.; Hunter, D.; Singh, A.; Vaia, R. Thermal degradation chemistry of alkyl quaternary ammonium montmorillonite. Chem. Mater., 13, 2979–2990 (2001) 5. Zhu, J.; Morgan, A. B.; Lamelas, F. J.; Wilkie, C. A. Fire properties of polystyrene–clay nanocomposites. Chem. Mater., 13, 3774–3780 (2001). 6. Gilman, J. W.; Awad, W. H.; Davis, R. D.; Shields, J.; Harris Jr., R. H.; Davis, C.; Morgan, A. B.; Sutto, T. E.; Callahan, J.; Trulove, P. C.; DeLong, H. C. Polymer/layered silicate nanocomposites from thermally stable trialkylimidazolium-treated montmorillonite. Chem. Mater., 14, 3776–3785 (2002). 7. Xie, W.; Xie, R.; Pan, W.-P.; Hunter, D.; Koene, B.; Tan, L.-S.; Vaia, R. Thermal stability of quaternary phosphonium modified montmorillonites. Chem. Mater. 4837–4845 (2002). 8. Davis, C. H.; Mathias, L. J.; Gilman, J. W.; Schiraldi, D. A.; Shields, J. R.; Trulove, P.; Sutto, T. E.; Delong, H. C. Effects of melt-processing conditions on the quality of poly (ethylene terephtalate) montmorillonite clay nanocomposites. J. Polym. Sci.: Part B: Polym. Phys., 40, 2661–2666 (2002). 9. Bottino, F. A.; Fabbri, E.; Fragala, I. L.; Malandrino, G.; Orestano, A.; Pilati, F.; Pollicino, A. Polystyrene–clay nanocomposites prepared with polymerizable imidazolium surfactants. Macromol. Rapid Commun., 24(18), 1079–1084 (2003) 10. Ray, S. S.; Okamoto, M. Polymer/layered silicate nanocomposites: a review from preparation to processing. Prog. Polym. Sci., 28(11), 1539–1641 (2003). 11. Awad, W. H.; Gilman, J. W.; Nyden, M.; Harris, R. H.; Sutto, T. E.; Callahan, J.; Trulove, P. C.; DeLong, H. C.; Fox, D. M.; Thermal degradation studies of alkyl-imidazolium salts and their application in nanocomposites. Thermochim. Acta, 409(1), 3–11 (2004).

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12. Davis, R. D.; Galman, J. W.; Sutto, T. W.; Callahan, J. H.; Trulove, P. C.; De Long, H. Improved thermal stability of organically modified layered silicates. Clays Clay Min., 52(2), 171–179 (2004). 13. Xiong, J.; Liu, Y.; Yang, X.; Wang, X. Thermal and mechanical properties of polyurethane/ montmorillonite nanocomposites based on a novel reactive modifier. Polym. Degrad. Stab., 86, 549–555 (2004). 14. Chigwada, G.; Jash, P.; Jiang, D. D.; Wilkie, C. A. Fire retardancy of vinyl ester nanocomposites: synergy with phosphorus-based fire retardants. Polym. Degrad. Stab., 89(1), 85–100 (2005). 15. Filho, F. G. R.; Melo, T. J. A.; Rabello, M. S.; Silva, S. M. L. Thermal stability of nanocomposites based on polypropylene and bentonite. Polym. Degrad. Stab., 89, 383–392 (2005). 16. Wang, S.; Zhang, Y.; Peng, Z.; Zhang, Y. New method for preparing polybutadiene rubber/ clay composites. J. Appl. Polym. Sci., 98, 227–237 (2005). 17. Burmistr, M. V.; Sukhyy, K. M.; Shilov, V. V.; Pissis, P.; Spanoudaki, A.; Sukha, I. V.; Tomilo, V. I.; Gomza, Y. P. Synthesis, structure, thermal and mechanical properties of nanocomposites based on linear polymers and layered silicates modified by polymeric quaternary ammonium salts (ionenes). Polymer, 46, 12226–12232 (2005). 18. He, H.; Duchet, J.; Galy, J.; Gerard, J. F. Influence of cationic surfactant removal on the thermal stability of organoclays. J. Coll. Interface Sci., 295(1), 202–208 (2006). 19. Hedley, C. B.; Yuan, G.; Theng, B. K. G. Thermal analysis of montmorillonites modified with quaternary phosphonium and ammonium surfactants. Appl. Clay Sci., 35, 180–188 (2007). 20. Kulkarni, P. S.; Branco, L. C.; Crespo, J. G.; Nunes, M. C.; Raymundo, A.; Afonso, C. A. M. Comparison of physicochemical properties of new ionic liquids based on imidazolium, quaternary ammonium, and guanidinium cations. Chem. Eur. J., 13, 8478–8488 (2007). 21. Cervantes-Uc, J. M.; Cauich-Rodriguez, J. V.; Vazquez-Torres, H.; Garfias-Mesias, L. F.; Paul, D. R. Thermal degradation of commercial available organoclays by TGA-FTIR. Thermochim. Acta, 457, 92–102 (2007). 22. Golebiewski, J.; Galeski, A. Thermal stability of nanoclay polypropylene composites by simultaneous DSC and TGA. Compos. Sci. Technol., 67, 3442–3447 (2007). 23. Araujo, E. M.; Barbosa, R.; Morais, C. R. S.; Soledade, L. E. B.; Souza, A. G.; Vieira, M. Q. Effects of organoclays on the thermal processing of PE/clay nanocomposites. J. Therm. Anal. Cal., 90(3), 841–848 (2007). 24. Costache, M. C.; Heidecker, M. J.; Manias, E.; Gupta, R. K.; Wilkie, C. A. Benzimidazolium surfactants for modification of clays for use with styrenic polymers. Polym. Degrad. Stab., 92(10), 1753–1762 (2007) 25. Modesti, M.; Besco, S.; Lorenzetti, A.; Zammarano, M.; Causin, V.; Marega, C.; Gilman, J. W.; Fox, D. M.; Trulove, P. C.; De Long, H. C.; Maupin, P. H. Imidazolium-modified clay-based ABS nanocomposites: a comparison between melt-blending and solutionsonication processes. Polym. Adv. Technol., 19, 1576–1583 (2008). 26. Calderon, J. U.; Lennox, B.; Kamal, M. R. Thermally stable phosphonium-montmorillonite organoclays. Appl. Clay Sci., 40(1–4), 90–98 (2008).

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SECTION II

PREPARATION AND CHARACTERIZATION OF RUBBER–CLAY NANOCOMPOSITES

CHAPTER 6

PROCESSING METHODS FOR THE PREPARATION OF RUBBER–CLAY NANOCOMPOSITES YONG-LAI LU LI-QUN ZHANG

6.1 INTRODUCTION According to the microstructure, the presently reported rubber–clay nanocomposites (RCNs) can be classified into three typical classes. (i) Intercalated: the silicate layers are partially expanded by rubber chains but the ordered structure is still retained. (ii) Exfoliated: the silicate layers are totally delaminated, and disorderedly dispersed in the rubber matrix. (iii) Separated: the nanosized silicate aggregates (not intercalated) are “separated” by the matrix rubber macromolecules, and usually the exfoliated silicates coexist in this type RCN. Among these three types of RCNs, the distribution homogeneity of the silicate layers in the exfoliated RCNs is highest and individual silicate layers act as reinforcing units, the aspect ratio of which are obviously larger than those of intercalated and separated silicate particles. Therefore, the exfoliated structure is regarded as the most ideal situation for RCNs in terms of properties enhancement. However, to achieve such exfoliated structure is quite difficult so that it is the most considered problem all the time in the researches in processing and preparation of RCNs. Moreover, a good interfacial interaction between the rubber and the nanoclay layers would enhance any effects on property improvement, most notably those of mechanical strength and gas-barrier capability. Besides, strengthening filler–rubber interaction would also favor good dispersion of silicate particles in the rubber matrix. As a result, the interfacial enhancement is another important problem for preparation of RCNs. Up to date, there are four types of processing methods having been developed for preparation of RCNs: (1) latex compounding, (2) melt intercalation, (3) solution intercalation, and (4) in situ polymerization intercalation. The former two methods Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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have more practical significance than the latter two. Therefore, this chapter will survey in detail the current situation relating to research into preparation of RCNs with latex compounding and melt compounding, including not only the mechanisms and influential factors linker to the morphological development of the RCN but also some methods used to enhance the interface between rubber and nano-layered clays. The works on solution intercalation and in situ polymerization will also be reviewed in brief. (Note: These two processing methods will be discussed into further details in the chapters of Section IV.) Finally, the future developing trend for processing method for preparation of RCNs will be analyzed from viewpoint of industrialization. 6.2 LATEX COMPOUNDING METHOD 6.2.1 Mechanism Layered pristine silicates can be dispersed in water, which acts as a swelling agent owing to hydration of the intercalary cations (usually Naþ ions). It should be noted that many rubbers are available in latex form, which is a more or less stable aqueous dispersion of fine rubber particles (in the submicron to micron range). During the late 1990s, Zhang et al. conducted a series of pioneering studies on the preparation of RCNs by mixing the rubber latex with a pristine clay water suspension, followed by coagulation [1–3]. The principle of this method was schematically represented in Figure 6.1. The clay layer could be steadily dispersed in water and the layers could be separated from each other. When the latex was added with vigorous stirring, the latex particle could be mixed uniformly with the layers; they interpenetrated each other. When clay–latex aqueous suspension mixture was coagulated together, the existence of latex particles was expected to prevent the dispersed clay layers from reaggregating, so that the dispersion state of clay layers in water could be kept and most of the layer could be dispersed in the rubber matrix on a nanometer level. Until now, many types of rubber-based clay nanocomposites, including natural rubber (NR) [4–9], styrene–butadiene rubber (SBR) [10–19], carboxylated nitrile butadiene rubber (XNBR) [11,16,20], nitrile butadiene rubber (NBR) [8,11,16,21,22], polyurethane (PU) [23], polybutylene rubber [24], and polychloroprene (CR) [6], have been successfully prepared via such latex compounding method (LCM). Compared with the melt, solution or in situ polymerization methods, LCM (when pristine

Clay

Water

Latex

Flocculant

Stirring violently

Mixing

Coagulating rapidly Nanocompound

FIGURE 6.1 A schematic of the formation of the nanocomposites through latex compounding method (LCM). Redrawn according to Figure 3 of Ref. [3] with permission from John Wiley & Sons.

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clay is employed) has demonstrated great promise for industrial application, due to the low cost of the pristine clay, the simplicity of the preparation process and the superior cost:performance ratio. 6.2.2 Influencing Factors The factors affecting the final dispersion level of clay in nanocomposites (NCs) prepared by LCM mainly include (i) the size of the rubber latex, (ii) the ratio of rubber latex to clay suspension, (iii) the processing approach of co-coagulation, and (iv) the interaction or affinity between rubber (latex particles) and clay layers. Apparently, the smaller the size of the latex particle and the higher the latex content, the better dispersion would be achieved. Nevertheless, these two factors could not be adjusted flexibly when the RCNs were prepared with LCM. Therefore, this chapter will only describe the latter two. 6.2.2.1 The Methods of Co-coagulation For co-coagulating the mixture of clay aqueous suspension and rubber latex, there are four processing routes having been reported, which would result in different dispersion state. In works [1–3,5,6,8,11,13,15,16,25] carried out by Zhang et al., several kinds of the electrolyte solutions, such as dilute hydrochloric acid solution, dilute triethylenetrammonium chloride solution, dilute calcium chloride solution, and dilute hydrochloric acid solution, were utilized as flocculants for co-coagulation of rubber latex/clay compounds. In the obtained nanocomposites based on different rubber latexes, there are coexisting both individual silicate layers and stacked silicate bundles with the thickness of about 10–30 nm, as shown in Figure 6.2a [16]. Moreover, the spatial distribution of clay in the rubber matrixes with different polarity is excellent. In most cases, however, the rubber macromolecules did not intercalate into the silicate galleries. As shown in Figure 6.3A, the positions of WAXD peaks corresponding to the regular stacking of silicate layers in RCNs prepared by LCM with the electrolyte flocculants are similar to those of Naþ -MMT after cation-exchange reactions with the flocculants (i.e., RNH3 þ -MMT and Ca2 þ MMT). Accordingly, the changes in WAXD diffraction peaks of rubber–clay nanocomposites by co-coagulation were expected to originate from the cations of flocculants in the interlayer space, rather than intercalation of rubber molecules. Therefore, these nanocomposites possessed a unique structure, referred to as a “separated” structure, where the rubber molecules “separated” the clay particles into either individual silicate layers or just layer aggregates of nanometer thickness without the intercalation of rubber molecules into clay interlayer spaces, different from intercalated and exfoliated clay nanocomposites. This separated structure resulted from the competition between separation of rubber latex particles and reaggregation of single silicate layers during the co-coagulating process with the electrolyte flocculants [16]. At the stage of mixing, the rubber latex particles were mixed with the clay aqueous suspension, in which clay was dispersed into individual silicate layers. After adding the electrolyte flocculants, the rubber latex and the silicate layers are coagulated simultaneously. The cations of the flocculants cause

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FIGURE 6.2 The microscopic (TEM and SEM) morphologies of rubber–pristine clay nanocomposites through LCM with different co-coagulation routes. (a) Using the electrolyte solution as flocculants (2% dilute triethylenetetrammonium chloride solution for the NR and SBR systems and 1% calcium chloride aqueous solution for the NBR and CNBR systems); the content of clay in these four types of RCNs is 20 phr. Reproduced from the Ref. [16] with permission from Elsevier. (b) NR/sodium fluorohectorite nanocomposite (100/20). The rubber latex/clay compound was cast in a in a mold built of glass plates, and then the casting was dried in air. Reproduced from Ref. [26] with permission from John Wiley & Sons, Ltd. (c) NR/PU (1:1) blend–clay nanocomposite (clay loading is 10 phrs.). The co-coagulation route is same as that for (b). Reproduced from Ref. [23] with permission from John Wiley & Sons, Ltd. (d) NR/ clay nanocomposites (100/5). The rubber latex/clay compound was concentrated in a spinning evaporator to remove water under reduced pressure. The produced thick slurry was cast on the flat plastic mold, and then dried in hot air (50–70 C). Reproduced from Ref. [7] with permission from Elsevier. (e) NR/clay aerogel nanocomposite (100/3). The water in the rubber latex/clay compound was removed by a process of freeze-drying. Reproduced from Ref. [9] with permission from Elsevier.

separated silicate layers to reaggregate so that the rubber latex particles between the silicate layers may be expelled. As a result, there are some nonexfoliated layers in the nanocomposites. In the meantime, due to the fact that the amount of latex is much more than that of silicate layers and the latex particles agglomerate rapidly, the reaggregation of silicate layers is evidently obstructed by the agglomerated latex

LATEX COMPOUNDING METHOD

1.52 nm 1.51 nm

5 phr

1.50 nm

10 phr

Intensity

(f) (e)

Intensity [a.u.]

(g)

1.31 nm

(d) 1.34 nm

(c)

20 phr

30 phr

1.34 nm

(b)

Pure clay

1.25 nm

(a) 2

151

6

4

2 3 4 5 6 7 8 9 10 11 12 13 14 15

10

8

2θ (º)

2θ (º)

(A)

(B)

800 700

Intensity

600

(c)

500 400 300

(b)

200 100

(a)

0 1

3

5

7

9

2θ (C)

FIGURE 6.3 WAXD patterns of Rubber–pristine clay nanocomposites by LCM with different co-coagulation routes. (A) Using the electrolyte solution as flocculants: (a) Naþ -MMT; (b) SBR–clay; (c) NR–clay; (d) RNH3þ -MMT; (e) NBR–clay; (f) CNBR–clay; (g) Ca2þ -MMT. The clay loadings for four RCNs are 20 phrs. The flocculant for SBR–clay, NR–clay is 2% dilute triethylenetetrammonium chloride solution (denoted as RNH3þ ), and that for NBR–clay and CNBR–clay is 1% calcium chloride aqueous solution (denoted as Ca2þ ). Reproduced from the Ref. [16] with permission from Elsevier. (B) NR/clay nanocomposites. The NR latex/clay compounds were concentrated in a spinning evaporator to remove water under reduced pressure. The produced thick slurries were cast on the flat plastic molds, and then dried in hot air (50–70 C) to form nanocomposites. Reproduced from Ref. [7] with permission from Elsevier. (C) NR/clay aerogel nanocomposites containing different amount of pristine MMT: (a) 1 phr; (b) 2 phr; (c) 3 phr. The water in the rubber latex/clay compound was removed by a process of freeze-drying. Reproduced from Ref. [9] with permission from Elsevier.

particles around the silicate layers. Consequently, the size of aggregates of silicate layers is at the nanometer level, and thus the obtained nanocomposites contain both the exfoliated silicate layers and the nonexfoliated (neither intercalated) aggregates of nanometer thickness in the rubber matrix.

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Some other researchers [4,7,22,23] performed the co-coagulation process by casting rubber composite film and air-drying, and achieved RCNs with different nanomorphologies from those prepared with the electrolyte flocculants. In the work conducted by Karger-Kocsis and coworkers [4], the NR latex was mixed with the aqueous dispersions of layered clays and the compounding ingredients (including vulcanization and antioxidant agents). The obtained compound latex was matured (keeping overnight to facilitate the diffusion of added chemicals), which ensures latex films to uniform properties, and then cast on raised glass plates having a dimension of 13 cm  10 cm  2 mm. The casting was then allowed to dry in air till becomes transparent and then vulcanized at 70 C for 4 h in an air circulated oven. The TEM image (as shown in Figure 6.2b) revealed that the intercalation/exfoliation of the silicates coexisted in the resultant RCN, and a skeleton (house of cards) silicate network was formed in the NR matrix. Afterwards, they prepared NR/PU (1:1) blend–clay nanocomposite with similar processing route [23]. The TEM image of the rubber–blend nanocomposite (see Figure 6.2c) displayed that NR and PUR are not compatible, and layered silicate stacks are mostly located at the boundary of the PUR (light) and NR (dark) phases. Pronounced intercalation and possible exfoliation took place only in the PUR phase. The silicate layers and aggregates cover the NR particles, resulting in a skeleton (house of cards) structure. It was suggested that the formation of this specific skeleton structure may yield improved mechanical properties. Galembeck and coworkers [7] employed a spinning evaporator (Rotavapor type) to first remove partial water from the clay/NR latex aqueous mixture under reduced pressure, and then cast the produced thick slurries on flat plastic molds, finally followed by hot air-drying at 50 or 70 C. The results of TEM (see Figure 6.2d) and WAXD (see Figure 6.3B) experiments indicated that well dispersion and almost full exfoliation of layered silicates could be achieved in the nanocomposite containing 5 phr clay. With clay loading increasing from 5 to 30 phr, the intercalated silicate structures become more dominate (see WAXD patterns in Figure 6.3B). Obviously, the speed of removing water from the rubber latex/clay aqueous compound during the co-coagulation process of casting film–air drying is far slower than that of using the electrolyte flocculants. Therefore, the silicate layers should have much more time to reaggregate during the course of air-drying. However, the rubber was not expelled from the silicate interlayer space, and nonintercalated stacked silicate bundles were not observed in the resultant RCNs. From above comparison, therefore, a conclusion can be drawn that the electrolyte flocculants should be essential reason for the formation of special “separated” structure in RCN prepared by LCM. The cations in the electrolyte have very strong interactions with the silicate layers having negative charges, which is far stronger than that between rubber latex particles and the silicate layers, resulting in rapidly reaggregation of silicate layers and excluding rubber from silicate interlayer spaces. Although the cocoagulation of film casting–air drying may result in the RCN with better dispersion, even full exfoliation, it is too time-consuming to be industrialized. Very recently, Lu and coworkers [27] reported a modified latex compounding method, where crosslinked rubber latex was used and the coagulation process was performed by spraydrying the mixture of cross-linked rubber latex and clay water suspension, instead of

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adding the flocculants. Consequently, in the rubber–clay compound obtained, the pristine clay layers were fully exfoliated. By incorporating this compound with other rubbers (using melt blending), it was possible to prepare rubber–pristine clay nanocomposites with an exfoliation morphology, even in EPDM. Pojanavaraphan and Magaraphan [9] also used freeze-drying technique to remove water from rubber latex/clay mixture and the produced nanocomposites exhibited the aerogel structure. The various amounts of prevulcanized NR latex and a 1 wt.% dispersion of pristine clay were mixed to prepare the nanocompounds. Then, the nanocompounds were immediately frozen within cylindrical glass shells at the temperature of liquid nitrogen. After freezing, the shells were attached to a lyophilizer (freeze-dryer) maintained at 54 C in a vacuum of less than 400 mTorr. The process for entirely subliming the ice out usually took 36 h. The SEM image (see Figure 6.2e) demonstrated that the clay aerogel structure was formed in some regions of the NR–clay nanocomposite containing 3 phr clay (i.e., NR/3MMT), and the observed aerogel structure was totally encapsulated with NR. The WAXD results (see Figure 6.3C) indicate that the intercalated structures existed in the nanocomposite at clay loadings of 1–2 phr, while the exfoliated structure was formed at clay loading of 3 phr. 6.2.2.2 The Interaction or Affinity Between Rubber and Clay It is naturally imaged that strong interaction or good affinity between rubber and clay would facilitate good dispersion of clay. He et al. [25] reported that a kind of full-exfoliated RCN (as proved by TEM image, see Figure 6.4a) could be prepared, in which 10 phr pristine Naþ -montmorillonite was dispersed in butadiene–styrene–vinylpyridine rubber (VPR) by latex compounding method followed by co-coagulation with acidic flocculants. In this research, the influence of the flocculant type on the microstructure was also studied. It was disclosed that the completely exfoliated structure was achieved using the Hþ flocculant (i.e., sulfuric acid, denoted as H-VPR/clay NC), whereas only the “separated” structure was obtained using the Ca2 þ flocculant (i.e., calcium chloride, denoted as Ca-VPR/clay NC). According to dynamic mechanical thermal analysis (DMTA) and solid-state 15N NMR characterization, it was suggested that a strong ionic interaction between the clay layers and the VPR molecules by means of Hþ , the pyridine group and the electronegative silicate layer, as schematically represented in Figure 6.4b. In consequence, H-VPR/clay NC exhibited obviously better mechanical and gas barrier performance than Ca-VPR/clay NC did. Unlike VPR, however, most rubbers are hydrophobic, so that their interfacial interactions with the hydrophilic pristine silicate will always be poor—a situation confirmed by the use of positron annihilation lifetime spectroscopy [28]. For this reason, most rubber–pristine clay nanocomposites prepared using LCM have one major shortcoming, namely that the modulus at high elongation (e.g., 300%) normally remains at a very low level, and so cannot meet the requirements of many engineering applications. In order to overcome this disadvantage, a number of studies [10,12,14,17–19,28,29] on interface enhancement with using organic modifiers have been conducted.

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PROCESSING METHODS FOR THE PREPARATION OF RUBBER–CLAY NANOCOMPOSITES

a Water

Latex

Stirring violently

Mixing

H+ Coagulating rapidly Nanocompound

b

Clay

b. VPR latex Na+ Na+

N

N

H+

.

H+

Na+ N

N

H+

a. Clay layers N+ H

Na+

Na+

N+ H

Na+

N

H+

Na+ Na+

Na+ Na+ Na+

Na+ Na+

+

Na+ Na+

Na+

Na+ Na+

Na+ Na+

Na+

N H

Na+ H+

H+

LATEX COMPOUNDING METHOD

155

Ma et al. [12] used three types of surfactant—triisopropanolamine (TA), m-xylylenediamine (MXD), and allylamine (AA) to first modify the pristine clay through cation-exchange reactions. The clay used to prepare RCNs via LCM was then replaced with surfactant-modified clays. As a result of this approach, exfoliation was achieved with the SBR/AA-modified clay, and enhancement of the mechanical properties was more prominent than when using carbon black. Similarly, Song and coworkers [29] used dimethyl ditollow-ammonium-modified clay to prepare SBR–NR–clay nanocomposites with LCM, and obtained the intercalated structure. Jia et al. [14] chose an unsaturated organic ammonium chloride (UOAC) to modify pristine clay in situ. For this, they added UOAC directly to a clay water suspension (before mixing with rubber latex) in the hope that, during vulcanization, a chemical interfacial interaction would form through reaction between the double bonds of UOAC and the rubber. In the nanocomposites thus obtained, an exfoliated, flocculant ions intercalated structure and a rubber intercalated structure were seen to coexist. The use of this approach was shown to improve the interfacial compatibility between clay and rubber, and to result in a major improvement in tensile strength. However, a strong chemical interface could not be formed, implied by that the modulus at high strain still remained at low level (Figure 6.5a). Jia and coworkers subsequently developed a relatively simple method for enhancing the interface, which included the addition of an organic modifier, hexadecyl ammonium bromide (C16) and 3-aminopropyltriethoxysilane (KH570), by melt blending after the latex compounding of rubber and pristine clay [17]. The results obtained (Figure 6.5b) showed that this approach could be used to tailor the interfacial interaction state to some extent by changing the modifier type and the flocculant system, and thus adjusting the shape of the strain–stress curve of the nanocomposites. Because the modulus at high strain was increased dramatically by the addition of KH570, it was assumed that a strong chemical interface might be formed. Afterward, a two-step method for interface enhancement was reported [18,28] in which KH570 was introduced to a clay water suspension before mixing with rubber latex; the bis(triethoxsilylpropyl)tetrasulfide (Si69) was then added by melt blending after latex compounding. The results obtained showed that with a combined modification of KH550 and Si69, the SBR–clay nanocomposites exhibited a maximum improvement in their mechanical properties (Figure 6.5c). The mechanical properties of SBR–clay nanocomposites with different interface enhancements (using optimum conditions for each approach) are summarized in Table 6.1. It can be seen from these data that the value of M300/M100 (which qualitatively may represent the interface state) for SBRCN with the two-step interface enhancement was highest. Based on these findings, Jia et al. proposed

3

FIGURE 6.4 High-resolution TEM image of butadiene–styrene–vinyl pyridine rubber (VPR)/clay nanocomposite (100/10) prepared through LCM with using 1 wt.% sulfuric acid solution as flocculant (left); and a schematic representation of an strong interaction between the pyridine groups of VPR and the electronegative silicate layers linked by Hþ (right). Reproduced from Ref. [25] with permission.

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PROCESSING METHODS FOR THE PREPARATION OF RUBBER–CLAY NANOCOMPOSITES

14 20

12

HC16

0.1

CaKH550

10

15

σ (MPa)

Stress (MPa)

HKH550

0.15

0.2

0.3 10

CaC16 8 6 Ca

H

4

5

Pure SBR

2 0 0

0 0

0

100 200 300 400 500 600 700 800

100 200 300 400 500 600 700 800

ε (%) (b)

Strain (%) (a) 18

SCK3S2

Stress (MPa)

15 SCK3

12 9

SCS2 6 SC 3

SBR

0 0

100

200

300

400 500

600

700

Strain (%) (c)

FIGURE 6.5 Strain–stress diagram of SBR/clay (100/10) nanocomposites prepared LCM. (a) Combing in situ modification with UOAC. 0, 0.1, 0.15, 0.2, and 0.3 represent the ratios of UAOC to MMT, respectively. Reproduced from Figure 9 of Ref. [14] with permission from John Wiley & Sons, Ltd. (b) Organic modifier was mixed by melt blending after latex compounding. C16 and KH550 denote hexadecyl trimethyl ammonium bromide, and 3-amionopropyl triethoxy silane, respectively; H and Ca represent Hþ and Caþ flocculated systems, respectively. Reproduced from Figure 7 of Ref. [17] with permission from John Wiley & Sons, Ltd. (c) Enhancing interface by a two-step method. SC, K, and S denote SBR–clay nanocomposite, KH550, and Si69 silane coupling agents, respectively. The number indicates the dosage (phr) for corresponding silanes. Reproduced from Figure 8a of Ref. [18] with permission from Elsevier. TABLE 6.1 Mechanical Properties of Interface Enhanced SBR/Clay (100/10) Nanocomposites by Different Approaches SBRCNs UAOC-0.2 HKH550 SCK3S2

M100 (MPa)

M300 (MPa)

M300/M100

Tensile Strength (MPa)

Elongation at Break (%)

Reference

1.5 1.8 2.3

2.5 6.8 9.6

1.67 3.78 4.17

18.7 12.4 16.9

638 504 509

[14] [17] [18]

Note: M100 and M300 represents modulus at 100 and 300% strain, respectively. The ratio of them (M300/ M100) was used to indicate the interface strength between filler and rubber.

MELT COMPOUNDING

157

the possible mechanism shown schematically in Figure 1 of Ref. [18]. Following co-coagulation, a KH550-modified (intercalated) nanoclay should be generated that contains many active hydroxyl groups originating from the hydrolysis of Si-(OR)3 of KH550 in aqueous suspension. During subsequent vulcanization, the ethoxysilylpropyl groups of Si69 would react with the Si–OH groups of nanodispersed clay in rubber, after which the sulfide groups of Si69 could react with the rubber molecules, assisted by accelerators of the curing system of rubber. As a result, a chemical bonding could be built between the clay phase and the rubber phase. Very Recently, Chakraborty et al. reported another two in situ modification techniques of pristine clay for preparing RCNs with LCM [10,19]. In these studies, naturally occurring unfractionated bentonite clay was used. One method is called in situ resol modification [10]. Predetermined amounts of resorcinol and formaldehyde (the molar ratio is 1/2) were directly added to the clay slurry, and reacted for 6 h under the condition of pH  8–9 and room temperature to produce resol modifiedbentonite clay. Incorporation of this modified clay in the SBR matrix by LCM with acid co-coagulation resulted in intercalation and/or exfoliation nanocomposites. The other method is called in situ sodium activation and organomodification [19]. The sodium chloride was added to the clay slurry for in situ sodium activation, and then, organomodification was carried out with assistant of octadecylamine (ODA). The in situ sodium activation was found to be a key step for this organomodification, and helped to increase the interlayer space distance from 1.28 to 1.88 nm. Incorporation of this modified clay in the SBR matrix by LCM also resulted in intercalation and/or exfoliation nanocomposites. In addition, a strong shear force during mixing of the curatives and other agents by melt blending after LCM would break down the silicate layers and clearly reduce the aspect ratio of dispersion of the RCN; this would be harmful to the material’s properties, notably those relating to the gas barrier. However, if the curatives are mixed in aqueous solution, and the rubber latex and clay water suspension mixture are used directly as a coating application, then this problem may be avoided. Indeed, the relative gas permeability (P/P0) of a thin coating RCN membrane has been reported to be as low as 0.02 when the clay content was 30 wt.% [24].

6.3 MELT COMPOUNDING 6.3.1 Mechanism Among the many currently available rubber compounding techniques, the melt compounding method (MCM) is probably the most widely applied for preparing RCNs, with related R&D studies having profited greatly from the general rules deduced from thermoplastic-based nanocomposites. The successful preparation of RCNs via melt compounding usually requires that the clay is first organically modified in order to (i) reduce its hydrophilicity and (ii) facilitate intercalation of the rubber chains into the gallery of clay layers.

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PROCESSING METHODS FOR THE PREPARATION OF RUBBER–CLAY NANOCOMPOSITES

Vaia and Giannelis were the first to establish the thermodynamic principles for polymer melt intercalation in organically modified clays (OMCs) [30,31]. Here, the free energy change (DF) upon intercalation was separated into an energy change (DE) due to new intermolecular interactions and a combinational entropy change (DS) associated with the conformational change of the polymer (DSpolymer) and the modifier’s aliphatic chains (DSchain); this could be expressed by following two equations: DF ¼ DE  TDS

ð6:1Þ

DS  DS chain þ DS polymer

ð6:2Þ

If DF is below zero, then the intercalation would be thermodynamically favorable; otherwise, it should be thermodynamically forbidden. For a nonpolar polymer, DE caused by the change in intercalated structures would be small, so that the intercalation should be dominantly determined by DS. According to both theoretical calculations [30] and experimental validation [31], there is a critical gallery height hc (2.4 nm; the corresponding basal spacing is 3.4 nm) for the compounding system containing clay modified by octadecylammonium salt (the initial basal spacing is 2.3 nm) and an arbitrary aliphatic polymer (nonpolar), below which the penalty for polymer confinement is compensated by entropy gains of the tethered surfactant chains associated with interlayer, and the overall entropy change is near zero. For a nonpolar polymer having weak interactions with surfactant chains and silicate surfaces, DF for intercalation of the polymer into the silicate gallery is always below zero if h < hc (where h is the gallery height of the layered silicates); in contrast, DF is larger than zero if h > hc. In some nonpolar rubbers, however, intercalation structures with a gallery height far larger than hc (i.e., 2.4 nm) were obtained by melt blending (see Figure 6.6, which shows the WAXD pattern and TEM image of the isoprene–isobutadiene rubber (IIR)/clay nanocompound). In order to interpret this phenomenon, the theory for melt intercalation was slightly modified by considering the special features of the rubber [32]. In comparison with common thermoplastic polymers, rubber has a much greater molecular weight, so that the viscosity and shear stress for rubber compounds during melt-blending remain quite high. Rubber chains are oriented by shear stress to a large extent, especially at the gap region between two rolls, and this may decrease the entropy of the chains before their intercalation. Therefore, DS polymer related to polymer confinement would be reduced during melt intercalation, so that hc for rubber in this dynamic state may be far larger than hc (2.4 nm) estimated based on thermoplastics. These analyses may also explain, from the viewpoint of thermodynamics, the influence of mixing shear magnitude during melt intercalation on the dispersion of OMCs. Moreover, the layered structure of silicates can be broken up upon high acting shear forces [33]. This effect is obviously the better the higher the molecular weight of the polymer (having higher melt viscosity). Rubbers are very high molecular weight materials showing extremely high viscosities during “melt” compounding, so

MELT COMPOUNDING

159

FIGURE 6.6 Microstructure and morphology of IIRCN (IIR/OMC ¼ 100/10) prepared by melt compounding: (a) WAXD patterns (the basal spacing of OMC is 2.4 nm.); (b) TEM images. Reproduced from Ref. [32] with permission from John Wiley & Sons Ltd.

that can supply much higher acting shear forces than thermoplastics do, facilitating silicate layer peel-off and dispersion in the matrix. Combining above considerations, a viewpoint has been brought out that rubbers should be the preferred matrices among polymers for melt intercalation of the silicate, and melt compounding should be a good way for achieving rubber nanocomposites [34,35], which has been validated by several researching results. For instance, it was reported by different research groups that directly mechanical (melt) compounding pristine clay (i.e., unorganically modified) with epoxidized natural rubber (ENR) could obtain intercalated nanocomposites [34,36]. Another sample is that melt compounding OMC with nonpolar matrix rubber, such as butadiene rubber (BR) [35,37], SBR [38], ethylene–propylene–diene–monomer rubber (EPDM) [39,40], isoprene–isobutyl rubber [41], without adding any other compatibilizers, can achieve nano-level dispersion of silicate particles at different extent. In contrast, via melt compounding without assistance of the compatibilzers, OMC is very difficult to be dispersed at nano level in polyolefin thermoplastics matrix, such as polyethylene (PE) and polypropylene (PP), whose chemical features are similar to those of EPDM and IIR. Furthermore, Lu et al. [42] reported that a series of highly filled RCNs based on EPDM, SBR, and epichlorohydrin rubber (ECO) could be prepared by melt blending with traditional rubber processing technique and the highly filled RCNs (clay content is up to 60 wt.%) still have intercalated silicate structures. As well to known, the nanocomposite having such high clay concentration is almost impossible to be prepared based on the thermoplastic polymer via melt compounding. Of course, a fact also should be considered

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PROCESSING METHODS FOR THE PREPARATION OF RUBBER–CLAY NANOCOMPOSITES

that present processing facilities for rubber melt compounding (i.e., inner mixer and two-roll mill) can provide and bear larger mixing torques than those for thermoplastics (i.e., screw extruder). 6.3.2 Influencing Factors Studies about preparation of RCNs with melt compounding were started later than those on thermoplastic–clay nanocomposites (TPCNs), and drew some lessons from them. As a result, some important factors determining the microstructure of the nanocomposite having been disclosed by studies on TPCNs had also been concerned by the researchers on RCNs. On the other hand, the key difference for rubber processing from that of thermoplastic is that vulcanization processing involving complex chemical reactions under high temperature and high pressure is necessary following melt compounding to obtain the final clay nanocomposites. Hence, the microstructural changes of RCNs during the vulcanization course and influence of OMC on vulcanization kinetic of matrix rubber had became important issues for melt compounding RCNs, attracting attention from many researchers. In this section, the factors determining microstructures of RCNs prepared with melt compounding will be introduced from two aspects: similar to TPCNs and unique for RCNs, as following. 6.3.2.1 Factors Similar to TPCNs Organic Modification of Clay At present, the organic ammonium compound is the mostly used intercalant for clay modification. The results of studies on NBR–OMC nanocomposites [43,44] have indicated that increasing the alkyl length of the surfactant would not only improve the spatial dispersion state (i.e., the dispersion dimension reduced but distribution homogeneity increased; see Ref. [43]) but also change the local structure in terms of intercalation and exfoliation (see Ref. [44]). Apart from the alkyl length of the surfactant, Zhang and coworkers also recognized that the intercalation extent increased with increasing surfactant dosage in the EPDM–OMC nanocomposites [45]. A study involving the SBR/OMC system also confirmed the above-described point that longer alkyl chains and a larger dosage of the modifier would facilitate the formation of a better dispersion state [46]. Zheng et al. [47] disclosed that increasing the polarity of the surfactant would facilitate the intercalation and delamination of OMC. Their results showed that the MMT modified with trimethyloctadecylamine or dimethylbenzyloctadecylamine (without polar groups) existed in the form of an intercalated layer structure, while the MMT modified with methyl-bis(2-hydroxyethyl)cocoalkylamine (containing polar groups) was fully exfoliated in the EPDM matrix. Besides the organic ammonium compound, some other surfactants were utilized to modify clay. For instance, Avalos et al. [48] used two quaternary phosphonium salts (aromatic and aliphatic) have been used as intercalants for Na-MMT and prepared NR nanocomposites with phosphonium salts modified clay via melt compounding method. It was found that the aliphatic salt was easier to intercalate into the clay galleries, due to its lower rigid structure, giving rise to a higher interlayer

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distance and facilitating the rubber intercalation obtaining an exfoliated structure in the nanocomposite. Liu et al. [49] reported that modification of the MMT was carried out with maleic anhydride (MA), which acts as the intercalation agent for MMT and the vulcanizing agent for EPDM matrix (peroxide curing system), as well as the compatibilizer for the EPDM and MMT phases. With regarding that a previous organic treatment of layered clay dramatically increases the prize of the final material limiting its application at industrial level, some researchers [50–55] developed a new processing route for melt compounding method without resorting to the organophilization route of pristine clay, in which the exchange reactions between pristine clays and the organic modifier were performed directly in rubber medium (named in situ organic modification). Wang et al. [50,51] prepared intercalated BR–clay nanocomposites with this new melt compounding method. BR, pristine clay, and intercalatant (dimethyl dihydrogenated tallow ammonium chloride (DDAC)) were directly mixed in a Haake rheometer at a given temperature, which resulted in intercalation of BR chains into clay interlayer space. Galimberti and coworkers used similar method to successfully prepared isoprene rubber (IR) and NR–clay nanocomposites with pristine clay [52–54]. (Note: Chapters 7 and 12 will also describe the above-mentioned new route for melt compounding method in more detail.) Carretero-Gonz
alez et al. [55] prepared NR–Na-MMT nanocomposites in one step by using poly(ethylene glycol) (PEG) as dispersing agent. NR and pristine Na-MMT was directly melt compounding with adding small amount of PEG, which acting as swelling agents favoring the intercalation of rubber chains into the silicate interlayer spaces and providing substantially improved clay dispersion, finally resulting in intercalated/exfoliated miscible hybrids. Feature of Rubber and Compatibilizers or Coupling Agents As with thermoplastic-based clay nanocomposites, it is relatively straightforward to obtain a good dispersion of clay particles in the rubber matrix with high polarity, examples being NBR [43,44,56,57], hydrogenated nitrile butadiene rubber (HNBR) [58], ENR [34], maleic anhydride-grafted (MAH-g)-EPR [59], MAH-g-IIR [60], CR [61]. Wu et al. [62] reported that NR and SBR (i.e., nonpolar but unsaturated rubber) could be more easily intercalated into the intergallery of OC during melt compounding than EPDM (nonpolar and saturated rubber). Ma et al. [63] also discovered that the ethylene content had little influence on the intercalation structure and dispersion of OMC in EPDM, nevertheless, the improvement in tensile strength of OC–EPDM nanocomposites with high ethylene contents (67–70%) was larger than that of OC–EPDM nanocomposites with low ethylene contents (52–52.5%). This was attributed to the strain-induced crystallization of PE segments and the orientation of clay layers in OC–EPDM nanocomposites with high ethylene contents. In this case, polar rubber could be used as the compatibilizer for improving the dispersion state of OMC, as well as for the interaction with nonpolar rubber. For instance, ENR [63,64] and MAH-g-EPDM [65–67] were used for NR/OMC and EPDM/OMC, respectively. Kim et al. [68] disclosed that the addition of a silane coupling agent could enhance the dispersion of OMC (C18-MMT) in NBR. Liu et al. [69] developed a reactive mixing

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intercalation method for preparation of NBR–OMC nanocomposites by introduction of the resorcinol and hexamethylenetetramine complex (RH) during melt compounding. RH acted as a reactive compatibilizer enhancing the interface combination between the rubber and the OMC through the interactions of RH with NBR and OMC. Processing Conditions Based on the results of their research into SBR/OMC and EPDM/OMC, both Sch€ on et al. [46] and Gatos et al. [67] drew the same conclusions, respectively, that a high shear mixing force could improve the dispersion of OMCs in a rubber matrix. Moreover, Gatos et al. [67] and Zheng et al. [40] demonstrated that a high temperature of compounding might even improve the dispersion of the clay and result in better mechanical properties. 6.3.2.2 Factors Unique for RCNs Influence of Mixing Curatives on Clay Dispersion In order to carry out vulcanization of the rubber, the curatives must be mixed into the compound. Some studies on EPDM/OMC [70,71] and IIR/OMC [71] indicated that mixing curatives into the compound would also change the intercalation structures. (Note: Both of two rubbers belong to nonpolar and saturated rubbers.) Goats and coworkers [70] observed further opening of the clay interlayer space after mixing curatives, and attributed this phenomenon to the “adsorption” of the curatives in the interlayer spaces. Ma et al. [71] studied the intercalated structures of EPDM or IIR/MMT modified by dimethyl dialkyl (C14–C18) ammonium (Nanomer I.44P) compounds at different mixing stage. The results demonstrated that the intercalation structure could not be directly achieved without other additives’ help, and steric acid (SA) should play key role on helping EPDM or IIR intercalation into OMC and improving clay dispersion. Evolutions of Microstructures of RCNs During Rubber Vulcanization The morphological change of RCNs during the vulcanization process can be classified as two types, namely changes in the local microstructure and in the spatial distribution of the clay particles. (i) Changes in the Local Microstructure of Clay Particles. This microstructural change could be deduced from making comparison between WAXD patterns of RCNs before and after curing, and can be further divided into three subclasses: . Further Intercalation. The position of the (001) reflection peak of the OMC is shifted to a lower 2u angle, or the peak may even disappear. This implies the formation of exfoliated structures or intercalated structures with a low degree of order. This phenomenon may caused by further intercalation of more rubber chains or low molecular mass substances (e.g., organic additives).

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.

Confinement or Collapse. The position of the (001) peak is obviously shifted to a higher angle after curing, but the basal spacing is still larger than that for a pure OMC used in this compound. This means that the originally interacted rubber chains moved outside the silicate interlayer space during the vulcanization process. . Deintercalation. The position of the (001) peak is larger than that of the pure OMC after curing. This suggests that not only the rubber chains moved outside the silicate gallery but also that there was deadsorption of the organic surfactant. In some cases, totally inorganic clays might even be formed. (ii) Change in the Spatial Distribution of Clay Particles. This type of change includes the dimension of dispersed clay particles and their distribution homogeneity on a large scale, with such information normally being acquired by using TEM. With regards to reinforcement with OMC, their spatial distribution homogeneity is more important than whether exfoliated clay layers exist. Hence, TEM observations at low magnification may be more suitable for detecting this type of morphological change. Recently, Lu and coworkers [72] observed the local microstructures and spatial distribution of OMC particles during the different stages of vulcanization of a variety of RCNs. Beyond all of Lu et al.’s and other research groups’ previous expectations, the results revealed that obvious microstructural changes at both local (Figure 6.7) and spatial scales (Figure 6.8) could occur over the entire vulcanization course, this being attributed to the high mobility of rubber segments in cross-linked network structures. The magnitude of change in morphology is most likely dependent on the matrix rubber polarity, with the morphological changes in nonpolar rubber-based nanocomposites (such as EPDM) being considerably greater than those in their polar rubber-based counterparts (such as BIIR). It should be noted here that the spatial dispersion state of clay particles in the EPDM–clay nanocomposite (EPDMCN) was improved to some extent during the early stage (i.e., scorch time, t10), but became far worse on completion of the entire vulcanization course (i.e., t100) (Figure 6.8). A number of studies have been conducted to identify the reasons and mechanisms for these morphological changes, the ultimate aim being to take advantage of this type changes in order to tailor the microstructures of the RCNs. At present, the study results have suggested that two classes of factors—namely chemical and physical— influence the morphological evolution of the RCNs during vulcanization. Chemical Factors. The chemical factors further include two aspects, notably (i) reactions between rubber and curatives at the initial stage of the curing process and (ii) reactions between curatives and amine-type intercalants within the silicate galleries. When Usuki et al. [73] investigated the influence of the curing accelerator type on microstructures of cross-linked EPDMCNs, the silicate layers of the clay were shown to be exfoliated and almost to disperse as

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FIGURE 6.7 Evolutions of local microstructures of EPDMCNs and BIIRCNs undergoing different stage of vulcanization: (a) WAXD patterns for EPDMCNs; (b) WAXD patterns for BIIRCNs; (c) d001 changing trends with curing course. t0 sample represents the uncured RCNs. t10, t20, . . ., t90, and t100 represents the time when the curing course reached 10, 20, . . ., 90, and 100% of full vulcanization course, respectively. Reproduced from Ref. [72] with permission.

monolayers in the cured EPDMCNs when a thiuram (TS) or a dithiocarbamate (PZ)-type vulcanization accelerator was used. It was also assumed that radicals produced by the thermal dissociation of TS or PZ could combine with carbon atoms in EPDM chains to polarize the EPDM molecules; this in turn would

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FIGURE 6.8 TEM images of (a) EPDMCNs and (b) BIIRCNs at different curing stage. Reproduced from Ref. [72] with permission.

result in a further intercalation of EPDM molecules into the clay interlayer spaces through hydrogen bonds between the polarized EPDM and the clay surface. Other studies of NR [74], ENR [34], EPDM [75,76], and HNBR [77]based clay nanocomposites, as conducted by Karger-Kocsis et al., revealed that the confinement, deintercalation, and further intercalation could occur simultaneously during vulcanization of the RCNs. Hence, these authors speculated that a Zn complex containing sulfur and amine groups from the organic amine-type intercalants might be formed during vulcanization. This Zn complex would serve as an intermediate for sulfur vulcanization, and react with rubber. However, if it were to migrate into the rubber matrix to take part in the vulcanization process, then the intercalants would be extracted from the clay galleries, with a resultant confinement and even deintercalation. But, if the complex were to cause rubber cross-linking inside the interlayer spaces, then the rubber molecules would be further inserted into the clay galleries, resulting in separation or delamination/exfoliation of silicate layers. It was further disclosed that the type of amine intercalant should play important roles in the structural changes of clay layers. For example, when a primary amine (e.g., octadecylamine) was used as the intercalant, then confined and even deintercalated structures were always found in the cross-linked RCNs. In contrast, when a quaternary amine (which has less reactivity with curatives) was used, then confinement and deintercalation was almost nonexistent [76,77]. When Karger-Kocsis and colleagues used peroxide as a substitute for sulfur curatives to cure HNBR–OMC mixing compounds, the RCNs obtained demonstrated a well-ordered intercalated structure, despite a primary amine having been used as the intercalant [78].

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Lu et al. reported a dramatic impact of curing temperature on microstructural changes in both the local and spatial scales of sulfur-cured IIRCNs [79]. Here, it was found that significant deintercalation occurred only when the curing temperature was above a certain level (i.e., 140 C in this system). Moreover, a WAXD investigation (Figure 4 in Ref. [79]) of the mixture containing OMC and sulfur curatives after thermal treatment at different temperatures suggested that the reactions between curatives and amine intercalants that occurred at high temperature resulted in only layer separation and delamination rather than deintercalation in the absence of rubber. This result suggested that the cross-linking reactions among rubber, curatives, and amine might be responsible to the deintercalation. In fact, a number of studies indicated that organic ammonium-modified clay could accelerate sulfur vulcanization of the matrix rubber, which will be introduced in detail in the next section. Thus, it can be expected that the curing rate of rubber intercalated into the silicate interlayer space might be higher than that of the rubber outside; consequently, more rubber could be driven to intercalate into the silicate interlayer space during vulcanization if the deadsorption of intercalants did not occur. Therefore, the hypothesis defining the role of the curing reaction in microstructural changes, as proposed by Karger-Kocsis and colleagues, might be more reasonable. Physical Factors. The two physical factors influencing microstructural changes among RCNs were identified as the high temperature and high pressure applied during the vulcanization process. As noted in Section 6.2.1, thermodynamically unstable intercalated structures can be produced by melt compounding; and these structures would persist at room temperature due to the high viscosity of the compound [32]. At high temperature, however, they would transform spontaneously to thermodynamically stable structures. High pressure was also found to enhance and accelerate this transformation. Similar investigation [80] was carried out on HNBR/OMC system. It was found that only 5 min thermal treatment (at 50–170 C) under atmospheric pressure (AP) can result in obviously changes in microstructure—the originally exfoliated organoclay dispersion in untreated HNBR/OMC compounds transformed into the coexisting of exfoliated, intercalated, and aggregated structures (containing partly and totally deintercalated structures). The treatment under high pressure at 20 C and long-time deposition (i.e., 30 days at 20 C and AP) also could cause such microstructural changes. Hence, an alternative mechanism for deintercalation phenomenon during vulcanization was proposed that the HNBR chains, which were inserted in the clay galleries in compounding, as well as the aliphatic chains of organic modified agents in the interlayer, can be partly pulled out from the clay galleries by the physical relaxation motion. Perhaps more importantly, high pressure was shown to be the critical factor causing the aggregation of originally well-dispersed OMC particles [81]. As shown in Figure 6.9, the dimension of the dispersed OMC particles was clearly

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IIRCN cured at 15 MPa × 180ºC 4.28 nm

2000 1500 1000 500

2.08 nm 1.35 nm

IIRCN cured at 15 MPa × 180ºC

0 2500

Uncured IIRCN 2000 1500

Intensity (cps)

1000 6.65 nm

500

3.72 nm

Uncured IIRCN

0 2500

SBRCN cured at 15 MPa × 150ºC 2000

5.05 nm

1500 1000 2.12 nm

500

1.39 nm

0 2500

SBRCN cured at 15 MPa × 150ºC Uncured SBRCN

2000 1500 1000

6.43 nm

500

3.27 nm 2.08 nm

0

Uncured SBRCN 1

2

3

4

5

6

7

8

9

10

2θ (º) (a) WAXD patterns

(b) TEM images

FIGURE 6.9 (a) WAXD patterns and (b) TEM images of cured and uncured IIRCN and SBRCN prepared by melting intercalation. Reproduced from Ref. [78] with permission from Wiley-VCH.

increased, and their spatial distribution homogeneity reduced dramatically after high-pressure vulcanization, when compared to uncured compounds. In contrast, the spatial dispersion state would be greatly improved after vulcanization under AP (see Figure 6.10).

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2500

IIRCN cured at AP × 180ºC

2000 1500

Intensity (cps)

1000 500

4.45 nm

0 2500

IIRCN cured at AP × 180ºC

SBRCN cured at AP × 150ºC 2000 1500 4.52 nm

1000 500

2.18 nm

0 1

2

3

4

5

6

7

2θ (º) (a) WAXD patterns

8

9

10

SBRCN cured at AP × 150ºC (b) TEM images

FIGURE 6.10 (a) WAXD patterns and (b) TEM images of IIRCN and SBRCN cured at atmospheric pressure. The vertical scales of the WAXD patterns are the same as that for Figure 7a in order to allow comparison. Reproduced from Ref. [78] with permission from Wiley-VCH.

Various methods to reduce the negative effect of high pressure have been considered in order to optimize the properties of RCNs, the most direct approach being to perform the vulcanization under a lower pressure. The data listed in Table 6.2 show that not only the mechanical but also the gas-barrier properties of IIRCN, when cured at 3 MPa pressure, are clearly superior to those of IIRCN cured at 15 MPa [82]. These results indicate the essential role of the clay dispersion state on the properties of the RCNs. Another approach would be to adjust the vulcanization temperature [79] and, indeed, results have shown that reactions during the initial curing period lead to a

TABLE 6.2 A Comparison of Properties Between IIRCNs Cured at 15 and 3 MPa Pressure Curing Pressure (MPa) Properties 

Shore A hardness ( ) Elastic modulus, G0 (MPa) Tensile strength (MPa) Elongation at break (%) Tear strength (kN/m) Relative permeability (Pc/Pp)a a

15

3

43 4.9 14.8 657 19 0.70

47 8.0 20.0 631 23 0.54

Pc and Pp are permeability of the composite and pure polymer, respectively.

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further intercalation of rubber chains into the clay galleries, thus improving the spatial dispersion of the clay particles. High pressure, in turn, leads to the aggregation of the clay layers to form larger clay agglomerates, although increasing the curing rate by raising the temperature can lead to a reduction in such aggregation. However, extreme curing temperatures (e.g., 160–180 C for this system) also resulted in deadsorption of intercalants and the formation of larger amounts of inorganic clays, thus weakening the interface interaction between clay and rubber. Hence, the optimum curing temperature range for obtaining a better dispersion was identified as approximately 140 C for the IIRCN system investigated. Because the cross-linking density, filler dispersion state and filler–rubber interactions are greatly influenced by the curing temperature, the gas-barrier properties and mechanical performance of IIRCN cured at different temperature would be quite different. Subsequently, it was found accidentally that the vaporization of residual solvents occurring under high-temperature treatment at AP could expand and delaminate the clay galleries, and that this would result in a major improvement in the spatial distribution of clay particles in the RCNs [83]. Enlightened by these results, a new strategy for the preparation of RCNs by using a melt-blending process was developed. For this, a certain amount of an organic solvent and OMC were first mixed to obtain a preswelled OMC (PSOMC) that could be incorporated into the rubber by melt blending. During the curing process, the solvents within the silicate galleries were vaporized, which in turn caused exfoliation and prevented any aggregation of the silicate layers. The RCNs prepared in this way showed a better dispersion morphology, and consequently their mechanical and gas-barrier properties were superior to those of RCNs prepared by traditional melt blending. As an example, the tensile strength and gas-barrier properties of IIRCN (100/5) created using this new method were 75 and 15% higher, respectively, than those of a counterpart prepared via traditional melt blending [84,85]. Influence of OMC on Vulcanization Course of the Matrix Rubber Vulcanization, which involves a huge amount of energy to convert an unformed plastic material to an elastic end product, is one of the most important processes in the rubber industry. Therefore, the influence of addition of OMC on rubber vulcanization kinetic had attracted the concern of the researchers in the field of RCNs. Most studies indicted that addition of OMC would considerably increase curing rate and crosslinking density of various rubber compounds, including NR [86,87], ENR [34], BR [88,89], SBR [90], EAR [91], NBR [92], and NBR/SBR blend [93], while addition of unmodified clay (i.e., pristine clay) would cause different effect on curing process for different systems, slightly accelerating, not influencing and depressing all having been observed. Furthermore, some studies [91,92] revealed that OMC could remarkably lowered the activation energy for curing reaction, and the autocatalytic model should be suitable for describing the cure behavior of the system. At present, above phenomena were usually traced to the accelerator role of the amine functionality of the intercalant used. Lopez-Manchado [87] observed that the octadeclyamine-modified clay exhibited further noticeable accelerator effect in the vulcanization rate of NR than the ODA itself, which implied a synergetic effect

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between the clay filler and the amine. Karger-Kocsis and coworkers [90] found that incorporation of OMC increased in cross-linking density of the matrix rubber, but resulted in high elongation at break. This contradiction was interpreted by a model where the silicate layers are intercalated/exfoliated and encapsulated by a highly cross-linked rubber layer and these “rubberized” platelets are embedded in the lightly cross-linked bulk. On the other hand, some studies [94,95] on EPDM/OMC systems shown that addition of OMC prolonged the optimum cure time and reduced the crosslinking density of the composites. 6.4 SOLUTION INTERCALATION AND IN SITU POLYMERIZATION INTERCALATION Solution intercalation is a very reliable method although less interesting for the praxis due to usage of organic solvent and not being environmental friendly. “Reliable” means that the silicate is reproducibly dispersed at nano level in the rubber after solvent (toluene, methylethylketone, N,N-dimethylacetamide, etc.) removal. It should be noted that organic modification of layered clay is also necessary for this method. In past, the solution technique has been used in preparation clay nanocomposites based on various rubbers, such as NR [96], polyisoprene (IR) [97,98], ENR [99], SBR [100–102], BR [101], NBR [101,103], EPDM [104], IIR [105], FKM [106–110], polyepichlorohydrin (ECO) [111,112], EVA [113,114], and HNBR [115]. In comparison with latex compounding, melt compounding, and solution intercalation, the literatures about preparation of the RCN through in situ polymerization are seldom. Liao et al. [116] prepared polybutadiene (PB), polyisoprene (PI), and SBR–organic montmorillonite (OMMT) nanocomposites were prepared by in situ anionic intercalation polymerization. The results showed that a certain extent of exfoliated rubber/OMMT could be achieved. The incorporation of OMMT obviously changed the microstructure feature of PB and PI: the concentrations of the 1,2-unit, 3,4-unit, and trans-1,4-unit increased dramatically with an increasing concentration of OMMT, and the concentration of the cis-1,4 structure decreased. In addition, the influence of OMC on molecular weight and molecular weight distribution of the rubber was observed. Zhu and Wool [117] prepared bio-based elastomer–clay nanocomposites through in situ radical polymerization of acrylated oleic methyl ester (AOME), which was synthesized from soybean oil, with presence of OMC. The morphology of the resultant RCN varies from fully exfoliated to intercalated structure with different types of clay and clay concentrations. 6.5 SUMMARY AND PROSPECT After over 10 years of development and investigation, rubber–clay nanocomposites gradually enter the age of application and commercialization. Because RCNs exhibit greatly decreased gas permeability, dramatically improved fatigue resistance,

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superior flame retardancy and aging resistance, and good mechanical properties, they are expected to be promising in application of many rubber product [118]. The maximum advantage for latex compounding method is superior cost:performance ratio due to not needing the organophilization route of layered clay. In addition, some techniques had been developed to tailor the interfacial interaction strength between nanoclay and rubber to satisfy different mechanical property requirements. The manufacturers in rubber industry are extremely sensitive to the cost of rubber products. Therefore, RCNs prepared by LCM might be first commercially applied at large scale. In China, a product line for NR/clay with an annual capability of 1000 ton and a product line for SBR/clay with an annual capability of 10,000 ton had be established with LCM technique in Hainan Province and Jilin Province, respectively. The produced nanocomposites have begun to be applied in the inner tube, liner, tread compound for out-road tire (ORT), and convey belt. [113] In the future, the morphological changes during postmechanical blending of other additives and vulcanization and their influence on the properties, which have been studied in the filed of RCNs prepared by melt blending, should be paid more attention too. The advantages for MCM represent its flexibility and using the traditional rubber processing facilities. A previous organic treatment of the layered silicate is usually required for MCM and dramatically increases the prize of the final material. Hence, the RCNs prepared by MCM might be more promising for application in special purpose rubber products, the cost of which is less regarded. Of course, the primary studies [50–55] on MCM with pristine clay (i.e., in situ organic modification route) had exhibited the potential in considerably reduction of the product cost. The main disadvantage for solution intercalation method is using of organic solvents, and not environmental-friendly. Some rubbers were synthesized in the solvent, for instance, BR, EPDM, and IIR. If the solution intercalation were carried out after accomplishment of polymerization and before removing of the solvent, the disadvantage of this method could be overcome. Recently, Exxon Mobile Chemical Company reported [119] their ongoing works in developing IIR–clay nanocomposites with above-mentioned method for tire liner application. Finally, it should be noted that improving dispersion of clay layers, especially achieving exfoliation, and enhancing interfacial interaction would also be invariable subjects for improving current processing methods or developing new methods for preparation of RCNs in the future. ACKNOWLEDGMENT We would like to thank Mr. Zhi-Qiang Xiu and Mr. Tuo Ji for their assistance in collection of the literatures and preparation of the reference list. REFERENCES 1. Zhang, L.-Q.; Wang, Y.-Z.; Yu, D.-S.; Wang, Y.-Q.; Sun, Z.-H.Chinese Patent ZL98101496.8 (1998).

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CHAPTER 7

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES MAURIZIO GALIMBERTI VALERIA ROSARIA CIPOLLETTI SIMONA GIUDICE

7.1 INTRODUCTION 7.1.1 Focus, Objective and Structure of Chapter 7 The present chapter is focused on the morphology of rubber–clay nanocomposites (RCN), a subject whose relevance is demonstrated by the efforts paid both in the academic and industrial worlds for interpreting RCN properties on the basis of their nanostructure, trying to establish a structure–property correlation. The availability of nanotechnologies, such as for example, the electron microscopy, allowed and increasingly allows the study of RCN morphology, a subject however investigated with the help of other analytical techniques, for example, X-ray diffraction (XRD), and covered in almost all the papers on RCN available in the scientific literature. Aim of this chapter is to collect the harvest of available data and comments, with the help of a red line for their organization and interpretation. The chapter moves from the summary of the state of the art, taking into account the first elaborations for nanocomposites morphology that were based on the study of thermoplastic materials and were then substantially adopted also for the elastomeric ones. In the discussion of the available data and results, attention is paid to tell apart experimental findings and interpretations. The type of layered clay and organic modifiers was selected as the red line for the elaboration of the available results; pristine clays are first considered followed by organically modified clays, as a function of the type of the modifier.

Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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7.1.2 X-Ray Diffraction Analysis for the Investigation of RCN X-ray diffraction analysis is a technique largely used for the study of RCN morphology, as it makes it possible to ascertain the crystalline structure of both clays (pristine or organically modified) and composites and, in particular, it allows to study the crystalline order (if any) in the direction perpendicular to the clay structural layers, a crucial aspect in the frame of the mechanisms proposed for the formation of RCN (see Sections 7.2 and 7.9). In fact, attention is paid on (00‘) clay peaks and, moving from these data, authors discuss the structure of the interlayer space and in particular the occurring of polymer chains intercalation. It is worth summarizing here the basic concepts at the basis of XRD pattern interpretation. A material containing a crystalline phase, bombarded with X-rays, gives rise to a diffraction pattern, a given reflection of the pattern being associated with evenly spaced planes present in the crystal. In the Bragg model of diffraction, indicating plane spacing and orientation with d and with three Miller indices (hk‘), respectively, X-rays scattered from adjacent planes produce a constructive interference when the angle u, between the plane and both the incident and reflected beams, and the plane spacing is related with the X-ray wavelength l through the following equation, known as the Bragg law: 2d sin u ¼ nl. In the case of layered silicates, (00‘) indices indicate the presence of a crystalline order in the direction perpendicular to the structural layers: the (001) reflection is due to the evenly placed layers and higher order reflections, for example, (002) and (003), originate from a regular arrangement of substances located in the interlayer space. These indices are largely reported in the RCN literature and thus in this chapter. An expansion of d interlayer spacing corresponds to a shift of (001) reflection toward lower 2u angle values.

7.2 BACKGROUND FOR THE REVIEW OF RCN MORPHOLOGY As anticipated in Section 7.1.1, for reviewing the morphological studies on RCN, the present chapter makes reference to the common knowledge about layered clays and polymer–clay nanocomposites. A brief summary is proposed here, having first as sources the reviews already published on polymer–clay nanocomposites [1–6]. Some of the summarized points are presented in more detail in the previous chapters of this book. 7.2.1 Cationic Clays Used for the Preparation of Rubber Nanocomposites Available papers on RCN are almost exclusively based on cationic clays montmorillonite, bentonite, fluorohectorite (FHT), vermiculite (VMT), attapulgite, and fibrillar silicates. Their main characteristics and the rubber they were blended with will be presented below in this chapter. As discussed in Chapters 1–4 of this book, to become compatible with the rubber matrix, pristine clays are modified with organophilic substituents, basically with ammonium cations. The structure of

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CH3 (a) C18H37

(b) T(H)

NH2

N+

CH3

CH3 T(H)

T(H) (c) CH3

N+

(d) CH3

CH2

N+

CH2CH(CH2)3CH3

CH3 C2H5

CH3 T(H) N+

(e) HOCH2CH2

T(H) (f)

CH2CH2OH

CH3

CH3

N+

CH3

T(H)

FIGURE 7.1 Structure of organic modifiers for a pristine clay: octadecylamine (a); quaternary ammoniums with the following substituents: three methyls and one talloyl (b); two methyls, one benzyl, and one talloyl (c); two methyls, one ethylhexyl, and one talloyl (d); one methyl, one (hydrogenated) talloyl and two hydroxyethyls (e); two methyls, two (hydrogenated) talloyls (f).

ammonium cations discussed in this chapter is shown in Figure 7.1, whereas in Figure 7.2 are shown the XRD patterns of a Na-montmorillonite: pristine (Figure 7.2a); modified with an ammonium cation bearing one methyl, one tallow group, and two (2-hydroxyethyl) (Figure 7.2b); modified with an ammonium cation bearing two methyls and two tallow groups (Figure 7.2c). 3.5 nm c

1.85 nm b

I (a.u.)

1.22 nm a 1.84 nm

1.22 nm

2

4

6

8

10

12

2θ (º)

FIGURE 7.2 XRD patterns of a Na-montmorillonite: (a) pristine (Cloisite); (b) modified with an ammonium cation bearing one methyl, one tallow, and two (2-hydroxyethyl) groups (Cloisite 30B); (c) modified with an ammonium cation bearing two methyls and two hydrogenated tallow groups (Dellite 67G).

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7.2.2 Multiscale Organization of Layered Clays Chapter 1 introduces the concept of multiscale organization of a clay mineral. In the present chapter, considering the upper level of the clay organization in the rubber matrix, items such as distribution and dispersion or aggregation of clay minerals, pristine, or organically modified, are discussed. As the most important feature of clay minerals is that they are layered, the analysis of the lowest level of clay organization is focused on examining the presence in the rubber matrix of individual clay layers or of stacking of them and the mechanisms proposed for their formation. 7.2.3 Clay Distribution and Dispersion It is worth specifying the meaning of terms such as distribution and dispersion that are often used (also synonymously) to indicate the level of homogeneity of claycontaining compounds. An efficient distribution leads to uniform rubber–clay composites, with even presence of clay particles, either agglomerated or as individual stacks or layers. An efficient clay (or filler in general) dispersion leads to the disruption of clay agglomerates and to clay incorporation in the elastomeric matrix, as layers or stacks of few lamellae. 7.2.4 Clay Modification: Intercalation of Low Molecular Mass Substances Coming to the lowest level of clay organization, one has to first consider the modification of a layered clay with lipophilic substituents, discussed in Chapters 2 and 4. As data available in the literature are almost exclusively based on cationic clays, the modification with organophilic cations (ammonium cations) is in particular discussed in the mentioned chapters and also in the present one. Said organophilic cations are known to intercalate in the space between two opposite layers. The intercalation of low molecular mass substances, considering also chemicals other than the organophilic cations, is one of the keys adopted in the literature and hence in this chapter for commenting data available in the scientific literature on RCN. 7.2.5 Types of Polymer–Clay Composites In almost any scientific paper dealing with polymer–clay nanocomposites, and hence in any chapter of this book, the reader may find words such as microcomposites, nanocomposites, flocculation, intercalation, delamination, exfoliation, words that describe the lowest level of clay organization. Some of them were already used in Section 7.1. Figure 7.2 is an elaboration of what reported by Giannelis et al. [1], Dubois and coworkers [2], and Ray and Okamoto [3] in the first general reviews on polymer–clay nanocomposites. This picture of composites’ types, formed from the interaction of layered silicates and polymers, was elaborated essentially on the basis of results collected in thermoplastic materials, as very few reports on RCN were available in those

BACKGROUND FOR THE REVIEW OF RCN MORPHOLOGY

185

days. Successive interpretations of RCN data were mostly based on this picture. The meaning of terms and related structures is explained as follows: .

.

. .

.

Microcomposite. Separated Phases (Figure 7.3a). The clay, either pristine or organically modified, is not efficiently dispersed in the rubber matrix; agglomerates are present, of micron size, made of stacks of layers or house of cards and the material properties are the ones traditional for a microcomposite. Nanocomposites (Figure 7.3b–d). According to the standard definition, they are composite materials characterized by the presence of dispersed particles whose size is in the “nanoscale,” defined as “having one or more dimensions of the order of 100 nm or less” [7]. In the case of clay nanocomposites, an efficient dispersion of clays is achieved; they are present in the polymer matrix either as individual layers or as stacks of few lamellae. The following terms refer more specifically to the clay organization in a nanocomposite. Exfoliated Nanocomposites (Figure 7.3b). Individual clay layers are separated in a continuous polymer phase, with the clay loading affecting their relative distance [3]. The terms exfoliation and delamination are synonymously adopted in the literature and this chapter prefers the former one (exfoliated clays), as suggested in Chapter 1. Intercalated Nanocomposites (Figure 7.3c). This expression refers to the presence of polymer chains in the space between two opposite layers. In the mentioned reviews, a minor extent of intercalation was assumed to occur, with preferentially extended polymer chains: “a single (and sometimes more than one) extended polymer chain” [2], “a few molecular layers of polymer” [3]. This placement of one/few layers of polymer chains was assumed in order to justify

Layered silicate

(a) Separated phases (microcomposite)

(b) Exfoliated (nanocomposites)

Polymer

(c) Intercalated (nanocomposites)

(d) Intercalated and flocculated (nanocomposites)

FIGURE 7.3 Composites’ types: (a) microcomposite, separated phases; (b) exfoliated nano-composites; (c) intercalated nanocomposites; (d) intercalated and flocculated nanocomposites.

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TABLE 7.1 Papers on RCN Considered in the Chapter, as a Function of the Type of Rubber Rubber NR IR ENR BR SBR (H)NBR IIR, BIMS EPDM

.

References [13–41] [42–44] [28,30,34,45,46] [47–58] [15,17,21,33,35,50,52,56,58–86] [21,43,50,52,62,63,65,67,68,85–99] [72,80,100,101] [35,43,65,102–115]

both (i) the somewhat limited expansion of the interlayer distance and (ii) the presence of a crystalline order. In fact, the intercalation was reported to occur with “a well-ordered multilayer morphology built up with alternating polymeric and inorganic layers” [2], “in a crystallographically regular fashion, regardless of the clay to polymer ratio” [3], as “the repetitive multilayer structure is well preserved, allowing the interlayer spacing to be determined” [2]. It was reported that both intercalated and exfoliated nanocomposites could be synthesized also with “high molecular weight polymers” [1]. Intercalated and Flocculated Nanocomposites (Figure 7.3d). In this case, polymer chains are intercalated between two opposite layers, that are “flocculated due to hydroxylated edge–edge interaction” [3].

7.2.6 Specific Literature on RCN This chapter takes into consideration not only the above-mentioned general reviews on polymer–clay composites but also the specific reviews published on RCN [8–12], as well as the large number of papers dedicated to particular rubbers, dealing with particular subjects. Table 7.1 indicates the references to papers on RCN considered in this chapter, as a function of the type of rubber.

7.3 RUBBER–CLAY NANOCOMPOSITES WITH PRISTINE CLAYS Morphology of RCN based on pristine clays, either from natural sources or synthetic, is discussed. The analysis of the multiscale clay organization is developed as follows: (i) Distribution and dispersion of the silicate in a rubber matrix. (ii) Analysis of the lowest level of clay organization, from the distance between two opposite layers to the presence of molecules in the interlayer space, if any.

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187

TABLE 7.2 Pristine Cationic Clays Used for the Preparation of RCN Clay

CEC (mequiv./100 g)

d Spacing (nm)

Na-Mt

76–125

1.10–1.30

Na-bentonite

59–80

1.24

Na-fluorohectorite

100

0.94

NR, IR, ENR, BR, SBR, (H)NBR, IIR, BIMS, EP(D)M NR, ENR, SBR, NBR, NR, ENR, SBR,

Rectorite

n.d.

2.15

SBR

Rubber

Mixing Technology Melt, solution, emulsion

Melt, solution, emulsion Melt, solution, emulsion Emulsion

7.3.1 Rubber Nanocomposites with Cationic Clays Table 7.2 reports type, cationic exchange capacity, and d spacing, as a range, of mostly used pristine cationic clays, and the rubber employed as the matrix. In almost all of the examined papers, clays were in the rubber in the absence of any further filler. 7.3.2 In a Nutshell In Table 7.3 the papers reporting data on RCN containing an unmodified silicate are organized. 7.3.2.1 Distribution and Dispersion of Pristine Clays in a Rubber Matrix As far as the distribution and dispersion of pristine clay in a rubber matrix are concerned, a summary can be attempted as follows: (i) The way to achieve an even distribution and dispersion of a hydrophilic pristine clay in a hydrophobic polymer matrix is to create a polar surrounding for both the clay and the polymer before mixing them. (ii) To mix clay and polymer dispersions separately prepared in polar fluids, preferably at a low clay content, appears as the right approach to overcome the mismatch of solubility parameter between clay and rubber and this holds for any type of layered silicates (bentonite, montmorillonite, fluorohectorite, as well as for fibrillar silicates). (iii) A processing method that foresees the use of a polymer latex appears as the ideal one to achieve an even clay dispersion; both the melt blending and the use of a polymer solution are not able to create the above-mentioned polar surrounding.

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TABLE 7.3 RCN Based on Pristine Clays Type of Clay

Type of Rubber

Blending Method

Na-Mt

NR

Melt Latex Solution Melt Solution Melt Melt Solution Melt

IR ENR BR SBR

(H)NBR

IIR, BIMS EPDM Na-bentonite

NR ENR SBR

References

Latex Solution Latex Solution Melt Melt Solution Melt

[13–16,27–30,39] [15,21–23] [20,116] [17,29] [44] [28,30] [47–49,53] [50] [15,17,50,59] (blend with NBR), [56,78] [21,71–73,77] [50,52,66–68,79] [15,21,85,87,88,98] [50,52,67,68,86] [50] (blend with SBR) [101] (IIR) [100] (BISM) [102,104,105]

Melt Latex Melt Melt Latex Solution

[19,31,32] [24,26,33] [45] [60] [33,74,75,77] [79] [19] [24–26,33] [45]

Na-fluorohectorite

NR ENR

Melt Latex Melt

Rectorite

SBR

Latex

[69,70]

Clay

NR BR SBR

Melt Melt Melt

[18] (calcined clay) [51] (china clay) [62] (clay powder)

Vermiculite

EPDM

Melt

[103]

Fibrillar silicate

SBR

Melt Latex Melt

[61] (attapulgite), [63] [61] (attapulgite), [64] [63,65]

(X)NBR

(iv) An improvement of clay dispersion seems to be also brought about, though to a minor extent, by the presence of polar groups on polymer chains. Results are in the following sections discussed as a function of the processing method as well as of the rubber. The best-achieved silicate dispersion is indicated in Table 7.4.

RUBBER–CLAY NANOCOMPOSITES WITH PRISTINE CLAYS

189

TABLE 7.4 Best Dispersion of a Pristine Clay in a Rubber Matrix as a Function of the Blending Method Blending Method Melt

Clay Mt, bentonite, FHT

Claya, vermiculitea, fibrillar silicatea

Rubber

Best Silicate Dispersion

Max D(d001 Spacing) (nm), Comment

0.2, absence of Simple NR, IR, ENR, polymer chains incorporation of BR, SBR, intercalation agglomerated (H)NBR, IIR, clays, poor BIMS, EPDM dispersion, poor interfacial adhesion BR, SBR, NBR, Uniform dispersion, EPDM reduction of aggregates, improved interfacial adhesion

Solution

Mt, bentonite

NR, BR, SBR, (H)NBR, BIMS,

Emulsion

NR, SBR, (X) Mt, bentonite, rectorite, FHT NBR, PUR

0.14, intercalation Conventional of polymer microcomposite, chains clay tactoids and agglomerates, big lumps of not exfoliated particles Excellent clay dispersion, individual and stacking layers, nanocomposites

0.2, flocculants intercalationb; 0.5–2, polymer intercalationc

a

Silicate modifiers: acrylate, silane, titanate, maleic anhydride (see “Chemically Modified Clays” section). Procedure (i). c Procedures (ii, iv, vi). b

7.3.2.2 Organization of Aggregated Pristine Clays in a Rubber Matrix As to the lowest level of clay organization, the following comments can be spent: (i) In most cases, when a crystalline order was maintained between aggregated Mt layers, the interlayer distance remained essentially unchanged; as a consequence, the intercalation of polymer chains was excluded. (ii) A minor expansion of interlayer distance was reported in some papers: in the presence of the same values, the intercalation of polymer chains was hypothesized or excluded. (iii) A major expansion of interlayer distance was obtained when mixing was carried out in the presence of other compound ingredients.

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(iv) An increase of disorder (broadening of reflections in the XRD pattern) was observed passing from the clay to the composite. 7.3.3 Distribution and Dispersion of a Pristine Clay in a Rubber Matrix 7.3.3.1 Rubber–Pristine Clay Composites from Melt Blending Composites with Montmorillonite or Bentonite as Pristine Clays .

.

.

NR and IR as the Rubber. Melt blending of Mt and NR led to a simple incorporation of agglomerated clay, as revealed by TEM analysis [13], to undispersed tactoids [14], and to the formation of microcomposites [15,16], with some large clay aggregates. A very poor dispersion of clay aggregates, though smaller with respect to those of the pristine clay, was as well observed in IR matrix [17], as shown in Figure 7.4. BR as the Rubber. In BR–Mt composites, SEM revealed clay particles agglomerates ranging from about 10 to 100 nm, whereas larger particles, from about 10 to 50 mm, were observed on the composite fracture surface, as well as many microvoids, originating from the breaking off of clay particles from the BR matrix, as a consequence of a poor interface adhesion [47–49]. E-SBR as the Rubber. The same morphology commented above for IR–Mt composites [17] was observed in E-SBR: presence of some large Mt aggregates, with the same stacked structure of the clay powder [15,59], as well as of Bentonite aggregates (observed by SEM) [60].

FIGURE 7.4 TEM micrograph of Mt in IR matrix. Composite prepared by melt blending (see Ref. [17]).

RUBBER–CLAY NANOCOMPOSITES WITH PRISTINE CLAYS

.

.

191

E-SBR/NBR Blend. Mt was observed in a blend to improve the compatibility between E-SBR and NBR phases (S ¼ 23.5 and N ¼ 32 as wt.%) [50]; SEM pictures revealed that NBR domains were separated from the continuous SBR phase in the absence of the filler whereas almost a single phase was obtained in the presence of Mt. EPDM as the Rubber. A nanocomposite structure was not achieved by mixing pristine Mt with a polymer containing a minor amount of a polar group, EPDMg-ma, TEM showed primary clay particles with, in the best case, good microscale dispersion [102].

Other Clays as Pristine Silicates Vacuoles were revealed (by SEM) on fracture surfaces of a composite based on a china clay in a BR matrix [51], indicating a poor matrix–filler adhesion. Large agglomeration of a calcined clay was observed in NRbased compounds [18]. Silicates with a High Aspect Ratio Fluorohectorite layers were not well dispersed both in a NR [19] and in an ENR [45] matrix. Silicate layers were only partially peeled off, by locally acting shear stress [45], even in the polar ENR. The high aspect ratio of stacked and distorted fluorohectorite layers was evident from TEM micrographs, for both the rubbers, and was correlated with the reinforcing effect brought about by this silicate [19]. Reagglomeration of purified attapulgite during drying process was observed, through TEM analysis, in a SBR matrix, giving rise to a fibrillous network structure [61]; the poor interfacial adhesion with the rubber matrix was indicated by the smooth and clear fracture surface of the composites. Chemically Modified Clays Modification of pristine silicates with chemicals other than ammonium cations affected to some extent the above-reported picture. Modification with either trimethylol propane triacrylate or triethoxy vinyl silane reduced the aggregates’ size of an Indian clay in E-SBR [62]. A uniform distribution in a BR matrix was obtained for a china clay modified with a titanate coupling agent [neopentyl (diallyl) oxy, trineodecanonyl titanate] [51], with enhanced polymer– filler adhesion and, as a consequence, less vacuoles with smaller size on the composites fractured surfaces. A uniform dispersion and a morphology similar to the one obtained with short microfibers was obtained by melt blending either SBR or NBR with nanosized fibrillar silicates modified with bis(3-triethoxysilylpropyl)tetrasulfide [63–65], with number and dispersion of nanofibrils found to increase with FS concentration, maintaining the anisotropy of oriented fibrils, thus suggesting a positive effect of the shear forces. Vermiculite modified with maleic anhydride was exfoliated and uniformly dispersed in an EPDM matrix [103], with an intimate interaction between the individual VMT layers and the rubber, as suggested by a not clear interface in SEM pictures.

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7.3.3.2 Rubber–Pristine Clay Composites from a Rubber Solution Rubber was first dissolved and then mixed under a vigorous stirring with a pristine Mt, used as such or previously dispersed in a polar liquid. A conventional microcomposite with typical clay tactoids was obtained from a toluene solution of a master based on NR and containing zinc oxide, sulfur, stearic acid, and a sulfenamide [20]. Dispersion problems were observed by mixing a Mt predispersed in either water or ethylalcohol and a rubber solution: BR [50], E-SBR [50,52,66,67], and BIMS [100] were dissolved in toluene, NBR [50,68] in chloroform, and composites were cured with either a peroxide (BR, E-SBR, NBR) or a phenolic resin (BIMS). For a Mt content at 4 wt.%, big lumps of not exfoliated particles were observed in an E-SBR (23 wt.% as styrene content) [67] with size in the range of 40–50 nm [66], though the pristine clay was able to improve the surface smoothness of an extrudate obtained from a capillary rheometer [52]. Agglomerated pristine Mt, not uniformly distributed, was shown in TEM micrographs of composites based on BIMS [100] and in composites based on an NBR having 34 wt.% as acrylonitrile content [50,68]. 7.3.3.3 Rubber–Pristine Clay Composites from a Rubber Latex The following main approaches are available in the open literature for performing the emulsion compounding. (i) Mixing of clay dispersion in water and of a rubber latex; coagulation with the help of a coagulating agent; cast of the dispersion; evaporation; drying; addition of further ingredients by melt mixing. (ii) As Procedure (i), but with the addition of further ingredients as a slurry prior to coagulation. (iii) As Procedure (i), without the use of a coagulating agent and with the addition of further ingredients by melt mixing. (iv) As Procedure (i), without the use of a coagulating agent and without the addition of further ingredients. (v) Mixing of clay dispersion in water, addition of rubber latex, and further ingredients prior to coagulation. (vi) With the use of a prevulcanized latex. Procedure (i): Use of a Coagulating Agent and Addition of Further Ingredients by Melt Mixing Zhang first introduced the emulsion compounding, coagulating in an electrolite solution a mixture obtained by vigorously mixing an aqueous clay suspension and a rubber latex. Clay layers became “separated” from each other, thanks to the large amount of rubber latex particles, and further compound ingredients were then added by melt mixing, typically at two-roll mill. A cation-type coagulating agent, triethylenetetrammonium chloride (2 wt.% solution), was used in Refs [21,69,70]: for Mt in NR, SBR, NBR, XNBR [21] and for rectorite in SBR [69,70]. Diluted sulfuric acid was used for preparing Mt-based composites, in Refs [71] and [72] with SBR as the matrix (Mt was modified with amines) as well as in Ref. [15] with SBR, NR, and NBR. One percent calcium chloride aqueous

RUBBER–CLAY NANOCOMPOSITES WITH PRISTINE CLAYS

water

latex

193

coagulated

FIGURE 7.5 Scheme for the formation of nanocomposites from emulsion compounding, depicted on the basis of what reported by Zhang [69].

solution or 2% sulfuric acid solution was used in Ref. [73] for SBR. For preparing bentonite-based composites, coagulation occurred in a dilute hydrochloric acid solution in Refs [74] and [75] with SBR and SVBR (V: vinylpiridine) as the matrices. A scheme for the formation of nanocomposites from emulsion compounding was proposed by Zhang as reported in Figure 7.5. An excellent Mt dispersion (20 phr) was achieved in NR, SBR, NBR, and XNBR [21], with both individual and stacking silicate layers with a thickness of about 10–30 nm. Nanocomposites and microcomposites were respectively obtained from emulsion and melt compounding [15]. A finer dispersion and a better interfacial strength was found for composites from sulfuric acid coagulation, that is, for layers without calcium [73]. Bentonite layer bundles with a thickness of about 4–10 and length of about 200–300 nm were observed in SBR and SVBR matrices, with a minor bundles aggregation at 40 clay phr rather than at 20 phr. Dispersion of rectorite layers at a nanometer scale was in SBR [69,70], with individual sheets and prevailingly bundles, with stack thickness lower than 100 nm. The difference of aspect ratios between rectorite and Mt (silicate length was about 1 mm and about 200–300 nm, respectively) was reduced by the rubber processing. The mixing procedure proposed by Zhang was then reported in more recent works. A uniform dispersion of Mt (2, 5 phr) in NBR matrix was shown in a composite prepared from a NBR latex coagulated with a dilute solution of dichloroacetic acid, with small aggregates with few clay layers at 10 phr as clay content [87]. Uniformly dispersed bentonite platelets were shown by TEM in E-SBR-based composites (5, 10 phr) by using diluted sulfuric acid as coagulating agent with a better dispersion at the highest clay content [76]. Mt aggregates were formed by coagulating with a diluted HCl solution [77] and a homogeneous composite was obtained only upon modifying the clay with N,N-dimethyldodecylamine. Procedure (ii). Use of Coagulating Agent and Addition of Further Ingredients as a Slurry Prior to Coagulation Single Mt layers and a few multilayer bundles were observed up to 5% by weight of silicate content in NBR-based composites prepared by pouring, drying, and curing a mixture from an NBR latex and an aqueous clay dispersion containing potassium hydroxide, an aromatic polyglycol ether and a sodium salt of methylene-bis-naphthalene sulfonic acid as electrolyte, nonionic emulsifier, and surfactant, respectively [88].

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Procedure (iii). No Coagulating Agent and Addition of Further Ingredients by Melt Mixing A more homogenous dispersion for the clay with respect to silica was observed through SEM analysis of fracture surfaces of composites prepared by mixing an aqueous clay suspension and the rubber latex [22], casting, evaporating, drying the mixture, and then adding sulfur and traditional activators and accelerators. Procedure (iv). No Coagulating Agent without the Addition of Further Ingredients by Melt Mixing Individual Mt layers coarsely aligned along the nanocomposite dry film plane and a significant number of few lamellae tactoids (thickness of about 3 nm) were, at 5 phr content, in a NR-based composite prepared by mixing the aqueous clay dispersion and the NR latex, reducing the mixture to a slurry by evaporation, casting on a mold, and finally drying [23]. Mainly tactoids were observed in the 30 phr sample. An unexpected rubber clay compatibility was shown by ESI-TEM, with absence of voids at the rubber–clay interface. Procedure (v). No Coagulating Agent and Addition of Further Ingredients Prior to Coagulation Bentonite clusters and individual sodium fluorohectorite layers still forming a skeleton structure were observed in NR, at 10 phr content, in a cured cast film obtained from an emulsion prepared by mixing an aqueous clay dispersion, a NR latex and compound ingredients [8,24]. Procedure (vi). Use of a Prevulcanized Latex Fluorohectorite stacks with an increased distance between opposite layers and (to a minor extent) exfoliated lamellae were observed in TEM micrographs of composites prepared by mixing a prevulcanized NR latex with a clay aqueous dispersion and, optionally, a PUR latex. Emulsions were filtered, cast in a mould, and dried, by finally postvulcanizing the latex at high temperature (100 C) [25]. Layers were located at the phase boundary of the NR/PUR blend, with silicate aggregates covering NR particles. More exfoliated structure was observed for fluorohectorite (2.5, 5, 7.5 phr), as a consequence of the higher aqueous swelling ability, with respect to bentonite that showed also stacks and, mainly, layers bundles [26]. 7.3.4 Organization of Aggregated Pristine Clays This section discusses the lowest level of the pristine clay organization, in particular investigates the crystalline order, if any, in the direction perpendicular to the structural layers. As in Section 7.3.2, results are organized as a function of the mixing technology and of the type of clay. 7.3.4.1 Rubber–Pristine Clay Composites from Melt Blending Montmorillonite .

NR as the Rubber. Mt interlayer distance was found to remain essentially unchanged [27–29] in final compounds, prepared through a multistep process,

RUBBER–CLAY NANOCOMPOSITES WITH PRISTINE CLAYS

.

.

.

.

195

by adding the clay (with CEC ranging from 70 to 93 mequiv./100 g) before or after other ingredients (curatives). Mt (001) peak was observed to become much broader in the composite pattern, remaining at the same 2u value and decreasing its intensity [29]; a so-called delamination mechanism was proposed: exfoliated clays are obtained by peeling off layers from clay stacks, thanks to the shear stress, without passing through the intercalation of polymer chains. Also in the presence of a slight Mt (CEC ¼ 119 mequiv./100 g) d001 expansion (from 1.22 to 1.38, 1.32, 1.31, and 1.25 nm for NR containing 1, 3, 5, and 10 phr of clay), it was commented that unmodified Mt particles were simply incorporated into the NR matrix, as they were hardly intercalated by the apolar rubber. NR/ENR Blends. When NR was used in blend with ENR (either ENR 25 or ENR 50 at 25 and 50 phr) [28], a more pronounced broadening of (001) peak was observed, with a minor interlayer expansion (0.09 nm), commented with the intercalation of a portion of elastomer chains. ENR as the Rubber. Two broad peaks in XRD patterns of Mt composites were based on ENR 25 or ENR 50, with ZnO, stearic acid, sulfur, and accelerators as further ingredients [30], corresponding to interplane distances which ranges from 3.0 to 3.2 nm and from 1.2 to 1.4 nm and interpreted as due to the clay with and without intercalated polymer chains, respectively. BR and E-SBR as the Rubbers. Mt interlayer distance was found to remain unchanged [47–49] or to increase to a very minor extent (less than 0.1 nm) [53] in BR-based compounds, prepared through a multistep process with different mixers, also adding further ingredients such as ZnO, sulfur, and a sulfenamide [48,49]. (001) peak was observed to become much larger in the composite pattern. A d001 expansion of 0.16 nm was observed in E-SBR-based composites, in the presence of ZnO and stearic acid, and commented with a slight rubber intercalation [78]. EPDM as the Rubber. The same Mt (CEC ¼ 110 mequiv./100 g) interlayer distance was detected in composites prepared in a multistep process containing curatives as further ingredients [104,105].

Bentonite The interlayer distance of the bentonite clay remained unchanged in composites based on NR [19,31,32], ENR [45], and SBR [60], exploring different clay CEC (mequiv./100 g), from 59 [60] to 70 [31] to 80 [19,32], adding the vulcanization ingredients before the clay [31] or after the clay [19,32,60]. Fluorohectorite Na-fluorohectorite was as well investigated in NR [19] and in ENR [45], analogously observing that the (001) clay peak did not change its position in the XRD pattern. Other Silicates Analogous conclusion was reported for composites prepared on a two-roll mill, based on an Indian clay in NR [32] or on a fibrillar silicate (attapulgite) in SBR [61].

196

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

7.3.4.2 Rubber–Pristine Mt Composites from a Rubber Solution Dispersion problems were reported in Section 7.3.2.2 for Mt-based composites prepared from a rubber solution and a clay dispersion in a polar liquid. Clay aggregates either did not show any (00‘) reflection in the XRD pattern or revealed an interlamellar distance very close to the one in the pristine clay. In fact, (001) Mt peaks were not detected in XRD patterns of composites with Na-Mt in BR [50,52], in E-SBR (15 and 40 wt.% as styrene content) [66], and in NBR with 19 wt.% as acrylonitrile content [50,52,68]. Interlayer distances of 1.23–1.24 nm were instead determined in the case of composites with NBR having 34 and 50 wt.% as acrylonitrile content [50,52,68], being the Mt interlayer distance in the 1.16–1.26 nm range. With BIMS as the rubber [100], a d001 spacing was reported to be at about 1.30 nm. A larger interlayer Mt distance (2.01–2.16 nm) was calculated only in the case of an E-SBR (23 wt.% as styrene content) [50,52,66,68,79], from a very small peak, and it was interpreted with polymer chains intercalation. In the same polymer, a K-Mt did not show any (001) peak. 7.3.4.3 Rubber–Pristine Clay Composites from Emulsion Compounding As reported in Section 7.3.2.3, a prevailing exfoliation of clay layers was obtained from the emulsion compounding. The order in the (less abundant) pillared layers is discussed as follows. Procedure (i). Use of a Coagulating Agent and Addition of Further Ingredients by Melt Mixing The interlayer distance was found to be larger in NR-based composites (1.34–1.51 nm) than in pristine Mt (1.25 nm) [21]; rubber intercalation was considered unreasonable, as these distances were lower than the one typical for an organoclay (larger than 1.7 nm). Basal spacings were also found to be dependent on the flocculant agent. Values as 1.43 and 1.51 nm were determined in composites based on E-SBR [73], using either sulfuric acid or aqueous calcium chloride, respectively. An interlayer distance of about 1.4 nm had been already reported [71] for Mt in E-SBR commenting the nonexfoliated clay structure. The intercalation of a single polymer layer was originally proposed [74] for Mt–E-SBR composites, commenting interlayer distances from 1.24 to 1.46 nm. In the case of rectorite in E-SBR [69], a slight change of the interlamellar space, from 2.15 nm (pristine silicate) to 2.4 nm, was attributed to some ion exchange of flocculant. A general comment was proposed by Zhang et al. [69], nanocomposites from emulsion compounding contain single silicate layers and pillared layers without macromolecules intercalated therein. Analogous comment was proposed by other authors for NBR/Na-MMT samples [87], and an expansion of about 0.2 nm of the interlayer distance was, attributed to the flocculant (dichloroacetic acid) absorption. Bentonite did not show any variation of the interlamellar spacing in an E-SBR composite [117]. Procedure (ii). Use of Coagulating Agent and Addition of Further Ingredients as a Slurry Prior to Coagulation Composites prepared following Procedure (ii), in particular by mixing a Na-Mt and a NBR in the presence of an electrolite, an

RUBBER–CLAY NANOCOMPOSITES WITH CLAYS MODIFIED

197

emulsifier, and a surfactant (see above) showed many peaks in the XRD patterns, as a function of the clay content, corresponding to interlayer distances of 1.65, 2.65, 3.23, 3.33, and 3.36 nm, with two peaks in the XRD pattern of the composite with the highest bentonite content (7.5 wt.%) [88]. Procedure (iv). No Coagulating Agent without the Addition of Further Ingredients by Melt Mixing Composites of Mt in NR [23] showed peaks in XRD pattern at lower 2u values with respect to pristine Mt. Position, broadness, and intensity appeared to be a function of the Mt content, the narrowest peak with the highest intensity was observed at a 2u value close to one of the pristine Mt for 30 phr Mt content, whereas a low angle broad almost undetectable halo was present in XRD pattern of the composite with 5 phr of bentonite. The interlayer distance was thus measured to move from 1.17 nm for pristine Mt to 1.59, 1.69, and 1.72 nm for composites with Mt content equal to 30, 20, and 10 phr, respectively. These findings were interpreted with the intercalation of rubber molecules within the lamellae stacks or tactoids. Procedure (vi). Use of a Prevulcanized Latex Basal spacing of bentonite was 1.24 nm and was calculated to move to 1.33 and 0.92 nm in the NR-based composites [26,33]. In the case of fluorohectorite, distances were 1.15 and 0.94 nm in the pristine silicate and at 1.28 nm in the NR-based composite. In XSBR [33], bentonite showed a d spacing of 1.46 and 0.93 nm and fluorohectorite of 1.46 and 1.29 nm and the higher polarity of XSBR was seen as responsible for the higher interlamellar distance. The intercalation of rubber chains into the silicate galleries was assumed to occur.

7.4 RUBBER–CLAY NANOCOMPOSITES WITH CLAYS MODIFIED WITH PRIMARY ALKENYLAMINES A primary alkenylamine with the alkenyl group having at least 12 carbon atoms (C12) was used to modify a clay and then to prepare rubber composites. In most cases, an octadecylamine (ODA) (C18 as the alkenyl group) was used as the clay modifier. 7.4.1 In a Nutshell Type of clay, rubber, and processing method used for the composites preparation are summarized in Table 7.5. Main characteristics of the organoclays are reported in Tables 7.6–7.11. Results are discussed as a function of the type of clay and rubber used as the composite matrix. They can be summarized as follows (i) RCN were prepared mainly through melt blending and, in some cases, by using a rubber solution, by adding, in most cases, further ingredients (e.g., curatives).

198

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

TABLE 7.5 RCN Based on Clay Modified with Primary Alkenylamines Type of Clay

Type of Rubber

Blending Method

Mt

NR

Melt Solution Melt Solution Melt Solution Solution Melt Solution Melt Solution Melt Solution Melt Melt Melt

BR SBR NBR (H)NBR BIMS EPDM Bentonite Fluorohectorite

NR SBR NR EPDM HNBR

References [13,19,27,32,46] [116] [53] [50,52] [78,81] [21,50,52,66,79,80] [50,52] [50,89,91] [80] [106–109,115] [35,110] [31] [79] [35] [111] [91]

(ii) Exfoliated high aspect ratio platelets, tactoids composed of variable number of platelets, and aggregates of tactoids were present in the final composites. (iii) The C18 alkenyl group promoted a better dispersion with respect to the shorter ones. (iv) Absence of (00‘) peaks in the XRD pattern of the final composites was verified when the clay with the lowest CEC or when a rubber solution were used. (v) In many cases well-resolved (00‘) peaks were observed, in particular three peaks, with interplane distances higher and lower than the interlayer distance of the starting organoclay. (vi) The expansion of interlayer distance was interpreted with polymer chain intercalation, whereas peaks with lower interplane distances were commented with either the deintercalation of the amine modifier and clay reaggregation or as higher order basal reflections due to the crystalline organoclay with intercalated polymer chains. 7.4.2 Composites with Montmorillonite and Bentonite 7.4.2.1 NR as the Rubber Better organoclay dispersion in the rubber matrix was observed with ODA as the clay modifier: with C12 as the nitrogen substituent, a higher number of tactoids aggregates was in fact observed [13]. A crystalline order for the organoclay in the composite was observed in most cases, as shown by XRD data in Table 7.6. (00‘) peaks were instead absent in XRD pattern of composites

199

b

Solution

Melt

Blending Method

Nanomer I30P from Nanocor. Best-defined peaks among the published ones.

2.69



119

a

1.76 2.1a 2.90, C18 1.67, C12

d001 (nm)

70 120a 119

Mt CEC (mequiv./100 g)

Organoclay

Further Ingredients

ZnO, SA, MBTS, S

ZnO, SA, MBTS, PBN, S ZnO, SA, CBS, S – –

– 3.0 (shoulder) 3.64b (3.6–4.0) 3.29–3.90

1

–

– 1.68 1.86b (1.8–1.9) 1.50–1.70

2

–

– 1.30 1.25b (1.25–1.31) 1.20–1.30

3

(00‘) Interplane Distances in the Composites

TABLE 7.6 RCN Based on NR and Mt Modified with Primary Alkenylamines

[116]

[27,31] [19,32] [13]

References

200

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

prepared either from a rubber solution [116] or by using organoclays with the lowest CEC (70 mequiv./100 g), such as bentonite [31] or Mt, that showed a very disordered structure already as a clay powder [27]. More in particular, in Ref. [27] the (001) peak was identified in the composite with pristine Mt and disappeared upon modifying the clay with ODA. In Ref. [116], a small broad peak was detected only at 10 phr as organoclay content and a broad peak (with the same interlayer distance as in the starting organoclay) at a lower clay content (7 phr) was observed with Mt modified with C12 alkenyl, confirming the lower tendency of this modifier to promote clay exfoliation. As mentioned above, most composites showed well-resolved XRD patterns, and interplane distances pretty similar to each other [13,19,32]. An example of the composite pattern, taken from Ref. [13], is shown in Figure 7.6. An enlargement of d001 spacing was interpreted with polymer chain intercalation. A larger expansion of interlayer distance was observed in the case of NR/ENR (5–10 phr) blends [34,46] and it was commented that ENR facilitates the intercalation of polymer chains, though not the exfoliation. Lower interplane distances were commented with confinement of the OM–clay with interlayer distance reduction [19], amine modifier deintercalation as a consequence of the reaction with chemicals such as zinc oxide [32], as the second-order diffraction peak (d002) of the intercalated organoclay (peak 2) [13], with reaggregation of silicate layers (peak 2), and with reformation of unmodified Mt (peak 3) [13] in the case of dodecylamine as Mt modifier. 7.4.2.2 BR and SBR as the Rubber As in the case of NR, peaks were either clearly detected or hardly observed (if any) in composites from melt or solution blending, respectively. In fact, when a rubber solution was used, SBR-based composites did not present (00‘) peaks in XRD patterns, as shown in Table 7.7, 25000 e

Intensity (cps)

20000

15000 d

10000

a c

5000

b 0 2

4

6 2θ (º)

8

10

FIGURE 7.6 XRD pattern of (a) ODA-MMT; (b) NR/1 ODA-MMT; (c) NR/3 ODA-MMT; (d) NR/5 ODA-MMT; and (e) NR/10 ODA-MMT nanocomposites. Reproduced from Ref. [13].

RUBBER–CLAY NANOCOMPOSITES WITH CLAYS MODIFIED

201

TABLE 7.7 RCN Based on SBR and Mt Modified with Primary Alkenylamine (00‘) Interplane Distances in the Composites

Organoclay Mt CEC (mequiv./ 100g)

d001 (nm)

90

2.24

Blending Method Melt

120a

90 a b

2.10

1.8

Solution

Further Ingredients

1

2

3

References

ZnO, SA, MBTS, S ZnO, SA, S, CBS, TMTD CB þ ZnO, SA, S, CBS, TMTD

2.94

1.60

–

[78]

–

1.35

–

[81]

2.52

1.35

–

DCUP

–

–

–

[50,52,66, 68,79]b, [80]

Nanomer I30P from Nanocor. Clay ¼ bentonite.

in a wide range of styrene content: 15 wt.% [66], 23 wt.% [50,52,66,68,79,80], 40 wt. % [66]. Two small peaks were in XRD patterns of BR-based composites [50,52], as shown in Table 7.8. By melt blending the organoclay with the BR rubber, a progressive expansion of the interlayer distance was observed in Ref. [53], adding the curatives, passing the TABLE 7.8 RCN Based on BR and Mt Modified with Primary Alkenylamine (00‘) Interplane Distances in the Composites

Organoclay Mt CEC (mequiv./ 100 g)

d001 (nm)

90

2.08

Blending Method

Melt

– a

1.8

Solution

Further Ingredients

1

2

3

References

– ZnO, SA, CBS, S, uncured ZnO, SA, CBS, S, cured

2.22 2.88

– 1.58

– –

[53]

3.5

1.6

–

DCUP

2.1

1.2a

–

From 2u values evaluated from the published XRD pattern.

[50,52]

202

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

composite at the two-roll mill and finally vulcanizing. Expansion of interlayer distance was as well observed in SBR-based composites from melt blending [78,81]. Lower interplane distances with respect to the organoclay d001 spacing were commented with a collapse of clay layers. Carbon black was reported to favor the organoclay exfoliation. 7.4.2.3 (H)NBR as the Rubber Melt Blending. Composites based on HNBR and Mt-ODA showed an expansion of interlayer organoclay distance, that remained unchanged after a peroxide curing [89] and vice versa decreased after sulfur curing [90,91], as shown in Table 7.9. These findings were interpreted with polymer chains intercalation and with amine modifier deintercalation due to its reactivity with sulfur-based systems. The formation of a complex between the amine modifier and the zinc-based active sulfurating species and the subsequent deintercalation was invoked to explain peaks corresponding to interlayer distances lower than in the pristine organoclay. Peaks detectable in the XRD pattern seemed to depend on the acrylonitrile (ACN) content for composites from a toluene rubber solution [50,52]. Intercalation of polymer chain was however hypothesized. 7.4.2.4 EPDM as the Rubber A nanocomposite structure was obtained with Mt-ODA in an EPDM, as such or as a 1:1 blend with an EPDM grafted with either maleic anhydride (EPDM-g-ma) or glicidyl methacylate (EPDM-g-gma) [106,107]. Polymer chains intercalation, clay reaggregation, and a monolayer arrangement of the intercalant were invoked to explain the different interplane distances reported in Table 7.10. The reactivity of the amine with vulcanization ingredients was hypothesized by observing the effect on the XRD pattern of different sequence of ingredient addition during the melt compounding. The organoclay interlayer distance was instead found to remain the same in composites (apparently) prepared in the absence of any further ingredient [108,109]. A clay modified with hexadecyl amine was used to prepare composites from a rubber solution [35,110]. Crystalline structures with several XRD reflections were observed when a blend of EPDM with EVA was used. Polymer chains intercalation, partial expulsion, and complete deintercalation of the ammonium cation were commented to interpret interplane distances shown in Table 7.10. 7.4.2.5 BIMS as the Rubber With Mt-ODA in BIMS as the rubber [80], (00‘) peaks were not detected in XRD pattern and this was attributed to polymer intercalation, followed by clay exfoliation. 7.4.3 Composites with Fluorohectorite Modified with a Primary Alkenylamine Fluorohectorite modified with ODA was compounded with NR, EPDM, E-SBR [111], and HNBR [91]. As shown by the data of Table 7.11, crystalline aggregates were present in any of the composites, along with high aspect ratio

203

2.10a

1.8

120a

–

b

Solution

Melt

Blending Method Further Ingredients

DCUP

ZnO, MgO, TAIC, TOTM ZMMBI, sDPA, DCUP, uncured ZnO, MgO, TAIC, TOTM ZMMBI, sDPA, DCUP, cured ZnO, MgO, TOTM, ZMMBI, sDPA, CBS, TMTD, S, uncured ZnO, MgO, TOTM, ZMMBI, sDPA, CBS, TMTD, S, cured

Nanomer I30P from Nanocor. From 2u values evaluated from the published XRD pattern.

2.10a

120a

a

d001 (nm)

Mt CEC (mequiv./ 100 g)

Organoclay

– 1.23, 1.28, ACN ¼ 34 –

– 1.64, ACN ¼ 50 –

–

–

1.70

3.10

–

3

2.11–2.37, ACN ¼ 19 –

1.9b

1.24b

1.87b

3.74

3.85

–

2 1.80b

3.60

1

(00‘) Interplane Distances in the Composites

TABLE 7.9 RCN Based on (H)NBR and Mt Modified with Primary Alkenylamine

References

[50,52], NBR

[90,91], HNBR

[89], HNBR

204

2.2b

2.19

1.96, C16 2.04, C16

119

119

76.4 85

b

Solution

Melt

Blending Method

Further Ingredients

DCUP DCUP

– – ZnO, SA, S, ZDEC, unvulcanized Vulcanized – ZnO, SA, S, ZDEC, unvulcanized Vulcanized ZnO, SA, S, TMTD

ZnO, SA, S, CBS, MBT, ZDEC

Nanomer I30P from Nanocor. From 2u values evaluated from the published XRD pattern.

2.10

120a

a

d001 (nm)

Mt CEC (mequiv./ 100 g)

Organoclay

1.60 1.70, EPDM/EVA 50/50 w/w

1.31, EPDM-g-ma – –, EPDM-g-ma

–, EPDM-g-ma 2.19 2.7, EPDM-g-ma – 3.58, EPDM/EVA 50/50 w/w

1.53 –, EPDM-g-ma –, EPDM-g-ma

3.30 3.15, EPDM-g-ma 3.15, EPDM-g-ma

2 1.51 1.29, EPDM-g-gma 1.29, EPDM-g-ma – – –

3.28 3.92, EPDM-g-gma –, EPDM-g-ma 2.2b 2.32 3.08

1

3

– 1.18, EPDM/EVA 50/50 w/w

–, EPDM-g-ma – –, EPDM-g-ma

– –, EPDM-g-ma –, EPDM-g-ma

– 0.94, EPDM-g-gma –, EPDM-g-ma – – –

(00‘) Interplane Distances in the Composites

TABLE 7.10 RCN Based on EP(D)M and Mt Modified with Primary Alkenylamine

[110] [35]

[109]

[108] [107]

[106]

References

205

Melt

Melt

Blending Method

ZnO, SA, ZDC, S

ZnO, MgO, TOTM, ZMMBI, sDPA, CBS, TMTD, S, uncured ZnO, MgO, TOTM, ZMMBI, sDPA, CBS, TMTD, S, cured

Further Ingredients

From 2u values evaluated from the published XRD pattern.

2.24

NR, SBR, EPDM –

a

2.0

d001 (nm)

HNBR 100

FHT CEC (mequiv./ 100 g)

Organoclay

5.0a, shoulder, NR 5.5a, SBR 4.6a, EPDM

3.04

3.54

1

2.1a, weak, NR –, SBR 2.1a, EPDM

1.77

1.8a

2

1.4a, NR 1.3a, SBR 1.3a, EPDM

–

–

3

(00‘) Interplane Distances in the Composites

TABLE 7.11 RCN Based on Synthetic Fluorohectorite Modified with Primary Alkenylamine

[35]

[91]

References

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

206

exfoliated platelets, prevailing in the case of NR and HNBR as the rubbers. Interplane distances (larger and smaller than in d001 of the starting organoclay) were, respectively, commented with intercalation of polymer chains and modifier deintercalation [91]. Composites in HNBR based on either Mt (see Section 7.4.2.3) or FHT modified with ODA showed very similar XRD patterns.

7.5 RUBBER–CLAY NANOCOMPOSITES WITH CLAYS MODIFIED WITH AN AMMONIUM CATION HAVING THREE METHYLS AND ONE LONG-CHAIN ALKENYL SUBSTITUENTS The structure of the clay modifier examined in this section is very close to the one discussed in Section 7.4; in fact, instead of primary amines, quaternary ammoniums with three methyl and one long-chain alkenyl groups are considered either octadecyl (C18) or hexadecyl (C16) alkenyl substituents were used. 7.5.1 In a Nutshell Papers published on RCN with Mt modified with this type of ammonium cation are indicated in Table 7.12. A report is also available on modified vermiculite in NR [36]. Data on the organoclays are reported in Tables 7.13–7.16. Morphology of RCN based on this type of organoclay can be summarized as follows: (i) An even dispersion of the organoclay in the rubber matrix was achieved in any of the composites, prepared by adding further ingredients (e.g., curatives). (ii) A fully exfoliated morphology was not achieved, the number of stacked layers being apparently higher in composites from a rubber solution or melt blending (also with a polar rubber such as HNBR) rather than in composites from emulsion compounding.

TABLE 7.12 Rubber Nanocomposites Based on Clay Modified with an Ammonium Cation Having Three Methyls and One Long-Chain Alkenyl Substituents Type of Clay

Type of Rubber

Blending Method

NR

(H)NBR EPDM

Solution Emulsion Solution Emulsion Melt Melt

[116] [37] [54] [73] [89,90] [102,104,105,112]

NR

Melt

[36]

Na-Mt BR

Vermiculite

References

RUBBER–CLAY NANOCOMPOSITES WITH CLAYS MODIFIED

207

(iii) (00‘) peaks were thus observed in the composites’ XRD patterns, that revealed in most cases several peaks and whose resolution (number and also definition of peaks) was in some cases better than for the starting organoclay. (iv) Shape and position of peaks appeared very similar in XRD patterns of composites based on different rubbers as well as on a clay modified with a primary alkenylamine (see Section 7.4). (v) Interplane distances were found, in some cases, to be the same in the organoclay and in the composite. (vi) In the majority of papers, it was reported that an expansion of the clay d001 spacing, commented with the intercalation of polymer chains and curatives and, at the same time, a reduction of the interplane distances, interpreted with deintercalation of the reaction product between curatives and clay modifier. (vii) This type of reaction was invoked to justify the supposed intercalation of an apolar polymer (such as EPDM) into the clay interlayer space. 7.5.2 Composites with Montmorillonite and Bentonite 7.5.2.1 NR as the Rubber A polar rubber favored a larger shift of organoclay basal spacing, as shown in Table 7.13, when a Mt modified with a C18-trimethylammonium cation was melt compounded with either NR or NR/ENR blends. XRD pattern in Figure 7.7 shows this finding and more reflection peaks for composites with ENR [34,46]. Essentially, the same interplane distances were instead observed in composites prepared from solution blending with Mt modified with ammonium cations based either on C16 or on C18 alkenyl substituents [116]; moving from primary amines to ammonium cations as Mt modifiers, the XRD pattern resolution significantly improved. A prevailingly exfoliated morphology with the presence of silicate stacks of several nanometers thickness, but without any detectable (00‘) peak in the XRD pattern, was obtained when a water suspension of C18-based organoclay was slowly added to a NR latex [37], removing water by freezing and melt compounding further ingredients at two-roll mill. White silicate stacks of several hundred nanometers with black areas attributed to NR were observed in SEM micrographs of a composite prepared by melt blending NR and vermiculite modified with a C18-containing ammonium cation [36]. (00‘) reflections were not detected in XRD pattern and interpretation was in favor of polymer chains intercalation. 7.5.2.2 BR as the Rubber Hydroxyl terminated BR oligomers were compounded with Mt (CEC ¼ 100 mequiv./100 g) modified with C18-trimethylammonium cation [54]. The interaction between terminal hydroxyl groups and clay layers was commented as crucial for layer exfoliation and from XRD analysis data, shown in Table 7.14, it was proposed that an intercalated/exfoliated structure a weak and broad hump was found at about 4.5 (as 2u value) at RT but not at a temperature (80 C), higher than the composite transition temperature.

208

4.38

Vermiculite, 98.6

b

Melt

Emulsion

Solution

Melt

Blending Method

ZnO, SA, S, 6PPD, MBT ZnO, SA, S, accelerator

ZnO, SA, MBTS, S

ZnO, SA, IPPD, CBS, S

Further Ingredients

Nanomer I28E from Nanocor. From 2u values evaluated from the published XRD pattern.

n.r.

Mt, 90

a

2.01, C18 1.93, C16

2.4, 1.6

d001 (nm)

Mt, 119

120

a

Clay, CEC (mequiv./ 100 g)

Organoclay

–

–

1.98 1.98

3.53, NR 3.93, NR/ENR 50

1

2

–

–

1.44 1.44

1.8 , NR 2.0b, NR/ENR 50

b

–

–

1.14 1.14

– 1.35b, NR/ENR 50

3

(00‘) Interplane Distances in the Composites

[36]

[37]

[116]

[34,46]

References

TABLE 7.13 RCN Based on NR and Clay Modified with an Ammonium Cation Having Three Methyls and One Long-Chain Alkenyl Substituents

RUBBER–CLAY NANOCOMPOSITES WITH CLAYS MODIFIED

Organoclay NR–ENR–Organoclay NR–Organoclay

d (001) = 3.93 nm

Intensity (a.u.)

209

d (001) = 3.53 nm d (001) = 2.39 nm

0

2

6

4

8

10

2θ (º)

FIGURE 7.7 XRD pattern of organoclay, NR/ENR 50/organoclay, and NR/organoclay composites (10 phr as organoclay content). Reproduced from Ref. [34].

7.5.2.3 SBR as the Rubber A homogenous clay dispersion without cavities at the Mt/rubber interface was observed in composites with Mt modified with C16trimethylammonium, prepared through emulsion compounding with Procedure (i) of Section 7.3.3.3, coagulating with sulfuric acid [73]. Layers aggregates were as well detected and their formation was attributed to the organic modifier. XRD patterns (for both cured and uncured compounds) revealed peaks, interpreted with SBR chains intercalation. When coagulation was carried out with an aqueous calcium solution, the bigger Ca2 þ ion led to lower d001 values. 7.5.2.4 HNBR as the Rubber A composite morphology very similar to the one obtained with ODA as the Mt modifier was reported with Mt modified with C18trimethylammonium [89,90]. Mostly intercalated clays were in TEM micrographs, the mechanism of peeling apart layers from clay stacks was hypothesized and variations of interplane distances, reported in Table 7.15, were attributed not only to the lower reactivity of the ammonium cation (with respect to a primary amine) with curatives, bringing to a lower intercalation but also, upon curing, to a lower modifier deintercalation. 7.5.2.5 EPDM as the Rubber The prevailing presence of clay stacks, evenly dispersed in the rubber matrix, was seen in melt blended EP(D)M-based composites; interplane distances are collected in Table 7.16. The same (hk‘) values were found for the organoclay and the composite when C16 was the alkenyl substituent, also by performing the mixing in the presence of

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

210

TABLE 7.14 RCN Based on Either BR or SBR and Clay Modified with an Ammonium Cation Having Three Methyls and One Long-Chain Alkenyl Substituents (00‘) Interplane Distances in the Composites

Organoclay Mt CEC (mequiv./ 100 g) BR 100

d001 (nm) –

Blending Method Solution

SBR 78

–, C16 in situa

a

Emulsion

Further Ingredients – ZnO, SA, S, MBTS, TMTD, IPPD, Hþflocculated system, uncured ZnO, SA, S, MBTS, TMTD, IPPD, Hþflocculated, system, cured ZnO, SA, S, MBTS, TMTD, IPPD, Ca2 þ flocculated system, uncured ZnO, SA, S, MBTS, TMTD, IPPD, Ca2 þ flocculated system, cured

1

2

3

1.95

–

–

4.37

2.08

1.43

References [54]

[73]

4.38

2.19

1.44

3.83

2.03

1.51

3.83

2.01

1.46

Clay modified, in situ, with an ammonium cation having a C16 chain as long chain alkenyl substituent.

EPDM-g-ma. An increase in the interlayer distance was instead reported in the case of C18 modifier, with a final exfoliation in the presence of EPDM-g-ma, and a C18 alkenyl group was considered to have the minimum length to have intercalation of polymer chains [104,105]. Basal spacing was reported to increase for a C18-based organoclay, as a consequence of polymer chain intercalation, then able to give rise to clay exfoliation [102]. The same morphology, with prevailingly not exfoliated clays, was obtained for EPDM composites with C18-based organoclay over wide

RUBBER–CLAY NANOCOMPOSITES WITH CLAYS MODIFIED

211

TABLE 7.15 RCN Based on HNBR and Clay Modified with an Ammonium Cation Having Three Methyls and One Long-Chain Alkenyl Substituents (00‘) Interplane Distances in the Composites

Organoclay Mt CEC (mequiv./ 100 g)

119

a b

a

d001 (nm)

2.50

Blending Method

Melt

Further Ingredients ZnO, MgO, TAIC, TOTM, ZMMBI, DCUP, unvulcanized ZnO, MgO, TAIC, TOTM, ZMMBI, DCUP, vulcanized ZnO, MgO, TOTM, ZMMBI, CBS, TMTD, S, unvulcanized ZnO, MgO, TOTM, ZMMBI, CBS, TMTD, S, vulcanized

1

2

3 b

3.20

1.60

3.98

2.0b

–

3.53

1.9b

–

3.87

1.96b

–

References

–

[89,90]

Nanomer I28E from Nanocor. From 2u values evaluated from the published XRD pattern.

TABLE 7.16 RCN Based on EPDM and Clay Modified with an Ammonium Cation Having Three Methyls and One Long-Chain Alkenyl Substituents (00‘) Interplane Distances in the Composites

Organoclay Mt CEC (mequiv./ 100 g) 110

– – a b

d001, d002 (nm) 1.8, 1.2, C16a 2.7, 1.3, C18b 1.97 2.3

Blending Method Melt

Further Ingredients ZnO, SA, S, CBS – ZnO, SA, S, ZDC

2

3

References

–

1

1.8

1.2

[104]

4.0 4.6 4.6

2.0 n.r. 2.0

1.4 n.r. 1.4

[104,105] [102] [112]

Clay modified with an ammonium cation having a C16 chain as long chain alkenyl substituent. Clay modified with an ammonium cation having a C18 chain as long chain alkenyl substituent.

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

212

ranges of EPDM characteristics such as ethylene content (from 52 to 70 wt.%), ENB content (from 11 to 26 as iodine value), viscosity (from 48 to 92 as Mooney units) [112]. The same interplane distances were detected, analogous to the ones reported above in Section 7.4.2.4 with ODA as the clay modifier. The reaction of the ammonium cation with curatives, in particular with zinc dialkyl dithiocarbamate was invoked by the authors to justify the increase in the interlayer distance. XRD analysis was performed on stretched samples, detecting an increase of the layers orientation.

7.6 RUBBER–CLAY NANOCOMPOSITES WITH MONTMORILLONITE MODIFIED WITH TWO SUBSTITUENTS LARGER THAN METHYL In this section, rubber composites are discussed based on a montmorillonite modified with an ammonium cation having, as substituents of the nitrogen atom, two methyl groups, one long-chain alkyl group, and one shorter chain group. 7.6.1 In a Nutshell Types of ammonium cation substituents and of rubbers employed for the composites preparation are summarized in Table 7.17. Results available in the cited references can be summarized as follows: (i) The in situ preparation of the organoclay was presented, both in the rubber melt and in a clay suspension. (ii) A homogeneous clay dispersion was revealed by TEM analysis in any of the composites, with coexistence of isolated layers and clay stacks, a fully exfoliated morphology was not achieved.

TABLE 7.17 Rubber Nanocomposites Based on Clay Modified with an Ammonium Cation Having Two Methyls and Two Substituents Larger than Methyl Type of Clay Na-Mt

Substituent of Ammonium Cationa Hydrogenated tallow, benzyl

Vinylbenzyl, octadecyl, vinylbenzyl, dodecyl Hydrogenated tallow, ethylhexyl Allyl, octadecyl

Type of Rubber

Blending Method

References

NR IR BR SBR SBR

Melt Melt Melt Melt Melt

[42] [42] [55] [82] [78]

IR NR E-SBR

Solution Solution Latex

[30] [20] [83]

RUBBER–CLAY NANOCOMPOSITES WITH MONTMORILLONITE MODIFIED

213

(iii) Several (00l) peaks were in most cases observed in XRD composites’ patterns, more and better resolved than for the starting organoclays, interplane distances higher and lower than d001spacing of starting organoclay were detected. (iv) In most cases, expansion of the interlayer distance was interpreted with polymer intercalation, whereas lower interplane distances, when commented, were attributed to collapse of the clay structure. (v) An alternative interpretation was proposed, only low molecular mass chemicals (the ammonium cation) were intercalated in the interlayer space, whose variation was attributed to reorganization of the ammonium cation substituents. Rubber was seen as a reaction medium for the in situ organoclay preparation and further XRD peaks with respect to (001) were commented as higher order basal reflections. 7.6.2 Hydrogenated Tallow and Benzyl Groups as Ammonium Cation Substituents Data on starting organoclays and composites are reported in Table 7.18. 7.6.2.1 Composites Based on Isoprene Rubber (NR, IR) The organoclay was formed in situ in Ref. [42], by reacting Mt and the ammonium cation in the rubber medium (both IR and NR). XRD patterns showed up to four peaks (see Table 7.18) and the determined d001 value was confirmed by measurements performed on TEM pictures. Number and shape of the peaks and the ratio among the interplanes distances led the authors to hypothesize the intercalation of only low molecular mass substances, the rubber acting only as the reaction medium. 7.6.2.2 Composites Based on Butadiene Rubber (BR, SBR) A similar interlayer distance of 3.8 and 4 nm, from XRD and TEM analysis, respectively, was determined for BR-based composites prepared through melt blending [55] and the intercalation of polymer chains was hypothesized. A d spacing of 1.44 nm was instead reported for the same organoclay blended with E-SBR 1502 [82]. Mt was modified with an ammonium cation containing two methyl and one vinylbenzyl groups and either an octadecyl or a dodecyl substituent and composites were prepared in E-SBR 1502 [104]. The interplane distances were interpreted with polymer intercalation and with the collapse of interlayer space. 7.6.3 Hydrogenated Tallow and Ethylhexyl Groups as Ammonium Cation Substituents 7.6.3.1 Composites Based on Isoprene Rubber (NR, IR) The pattern with the highest number of well-defined peaks (see Figure 7.8 and Table 7.19) for composites with an organoclay with only one long-chain substituent was

214

b

a

Melt Melt

Melt

–a

1.9b

–b 2.32 vinylbenzyl octadecyl

2.05 vinylbenzyl dodecyl

Melt

Blending Method

–a

d001 (nm)

Organoclay prepared in situ in the rubber medium. Cloisite 10A from Southern Clay.

125 90

NR 128 IR 128 BR 125 SBR

CEC (mequiv./100 g)

Organoclay

ZnO, SA, S, MBTS, agerite, CDDC

ZnO, SA, CBS, S, IPPD

ZnO, SA, S, MBTS

–

–

Further Ingredients

3.15

– 3.15

3.8

4.02

–

1

–

– 1.73

–

1.96

2.15

2

1.2

1.44 1.2

–

1.32

1.4

3

1.0

–

4

(00‘) Interplane Distances in the Composites

TABLE 7.18 RCN Based on Mt Modified with Hydrogenated Tallow and Benzyl Group as Ammonium Cation Substituent

[82] [78]

[55]

[42]

[42]

References

RUBBER COMPOSITES WITH MONTMORILLONITE MODIFIED WITH AN AMMONIUM

215

1.104 d001 = 40.1 Å Natural rubber + 10 phr M3

Intensity (CPS)

8000

6000

d001 = 34.0 Å Pure M3

4000

2000

0 0

2

4

6

8

10

12

14

16

2θ (º)

FIGURE 7.8 XRD pattern of organoclay, dimethyl hydrogenated tallow (2-ethylhexyl) ammonium montmorillonite (M3), and polymer composites. Reproduced from Ref. [20].

obtained [20] for a composite prepared by mixing in toluene the organoclay and a rubber blend with further ingredients. As the largest interplane distance was 4.01 nm, the intercalation of polymer chains was commented. Similar interplane distances were reported in Ref. [30], upon mixing IR and organoclay in toluene and interlayer distance expansion was commented with polymer chain intercalation. 7.6.4 Other Long- and Short-Chain Alkenyl Groups as Ammonium Cation Substituents The organoclay was prepared in situ [83] by reacting Mt and an unsaturated organic ammonium salt, N-allyl-N,N-dimethyl-octadecylammonium chloride, then blended with E-SBR following the Procedure (i) reported in Section 7.3.3.3. The interlayer distance of pristine Mt (1.25 nm) was found to increase with the modifier/clay ratio, from 0.05 (1.34 nm) to 0.3 (1.9 nm). Interplane distances in Table 7.19 refer to this latter ratio, commented with the intercalation of polymer chains.

7.7 RUBBER COMPOSITES WITH MONTMORILLONITE MODIFIED WITH AN AMMONIUM CATION CONTAINING A POLAR GROUP Montmorillonite modified with methyl-tallow-bis-(2-hydroxyethyl)quaternary ammonium cation was used to prepare composites with different rubbers, in most cases through melt blending.

216

From 2u values evaluated from the published XRD pattern.

Latex

1.90, allyla

78

a

Solution

Blending Method

3.4, ethylhexyla 2.86–3.02, ethylhexyla

d001 (nm)

– –

Mt CEC (mequiv./ 100 g)

Organoclay

SBR ZnO, SA, S, DPG, MBTS, TMTD, IPPD

NR ZnO, SA, S, CBS ZnO, SA, S, calcium stearate, MBT, CBS

Further Ingredients

4.7

4.01 3.5–3.6

1

1.9, shoulder

2.0a 1.8a

2

1.4

1.33a 1.2a

3

–

1.0a –

4

(00‘) Interplane Distances in the Composites

[83]

[20] [30]

References

TABLE 7.19 RCN Based on Mt Modified with Hydrogenated Tallow and Ethylhexyl or Short-Chain Alkenyl Groups as Ammonium Cation Substituents

RUBBER COMPOSITES WITH MONTMORILLONITE MODIFIED WITH AN AMMONIUM

217

TABLE 7.20 Rubber Composites with Montmorillonite Modified with Ammonium Cation Containing a Polar Group (Methyltallow-bis-(2hydroxyethyl) Quaternary Ammonium Cation) Type of Rubber

Blending Method

References

NR

Melt Solution Melt Melt Melt Melt

[19,28,32] [20,30] [55] [84] [92] [89,90]

BR SBR NBR HNBR

7.7.1 In a Nutshell Type of employed rubber, processing method and papers available in the literature are indicated in Table 7.20. Characteristics of the organoclay are shown in Table 7.21. Results available in the open literature can be summarized as follows: (i) The presence of the polar hydroxyl group as ammonium cation substituent seems to cause a worse dispersion of the organoclay in the rubber matrix, with a larger amount of layers agglomeration. (ii) With this type of ammonium cation, crystalline organoclay structures are not present in the final composite; XRD patterns show only broad peaks, in most cases only one peak. (iii) Expansion of interlayer distances from starting organoclay to composites is reported and is interpreted with the intercalation of polymer chains. (iv) Interplane distances were lower with respect to d001 spacing of the organoclay, when detected were commented with deintercalation of clay modifier. (v) The functional group on the clay modifier could be used to promote a reaction with a functionalized polymer chain and expand the interlayer distance. 7.7.2 Composites with Diene Rubbers 7.7.2.1 NR and BR as the Rubber Less exfoliation and more agglomerated particles were seen with methyl tallow bis-2-hydroxyethyl ammonium modified Mt with respect to octadecylamine modified Mt [32], with properties resembling those of microcomposites, whereas a finely intercalated structure was seen only with Mt modified with dimethyl-dihydrogenated talloyl ammonium [28]. Expansion of interlayer distance calculated from XRD patterns was commented with intercalation of polymer chains. Peaks were similar to those of composites with Mt-ODA [32] and broader than in the starting organoclay [28]. A lower interplane distance with respect to organoclay powder [19,32] was considered as due deintercalation of clay modifier. 7.7.2.2 SBR as the Rubber In Ref. [84], to promote the clay exfoliation, a further modification of the organoclay was carried out, by reacting the hydroxyl

218

b

a

1.80a

Melt

1.85a

90

Melt

1.8a

MgO, ZnO, TOTM, ZMB, ZMMBI, sDPA, CBS, TMTD, S, uncured MgO, ZnO, TOTM, ZMB, ZMMBI, sDPA, CBS, TMTD, S, cured MgO, ZnO, TAIC,TOTM, ZMB, ZMMBI, sDPA, DCUP, uncured MgO, ZnO, TAIC,TOTM, ZMB, ZMMBI, sDPA, DCUP, cured

ZnO, SA, S, MBTS

ZnO, SA, S, MBTS

ZnO, SA, S, CBS ZnO, SA, S, MBTS, PBN ZnO, SA, S, Ca-stearate, MBTS, CBS

Further Ingredients

Cloisite 30B from Southern Clay. From 2u values evaluated from the published XRD pattern.

Melt

Solution

Melt

Blending Method

1.85a 1.85a 2.3

1.80a

d001 (nm)

NR 90 90 – BR 90 NBR (ACN ¼ 33wt.%) 90 HNBR (ACN ¼ 43wt.%) 90

Mt CEC (mequiv./100 g)

Organoclay

4.26

–

3.55

–

n.r.

4.2

3.0, shoulder 2.45 3.1

1

– 1.35b

2.1b

–

–

n.r.

–

1.30 – –

3

–

1.78b

–

n.r.

–

– – –

2

(00‘) Interplane Distances in the Composites

TABLE 7.21 RCN Based on Mt Modified with Hydrogenated Tallow and bis(2-hydoxyethyl) Methyl as Ammonium Cation

[89]

[90]

[92]

[55]

[19,32] [28] [30]

References

RUBBER NANOCOMPOSITES WITH MONTMORILLONITE MODIFIED

219

groups of the ammonium cation with maleinized low molecular mass BR; d spacing was found to move from 1.8 to about 7 nm and even to 14.0 nm, for BR with 7.5% and 2.5% maleic anhydride content, respectively. BR modifier with the lower maleic anhydride content was found to promote a better clay dispersion. 7.7.2.3 ENR as the Rubber In Ref. [30], mixing of ENR 50 dissolved in MEK and organoclay swollen in toluene was performed and an expansion of the interlayer distance was calculated from the broad peak in the XRD pattern, commented with polymer chains intercalation. 7.7.2.4 (H)NBR as the Rubber Ref. [92] suggests that the formation of intercalated nanocomposites is favored by the presence of polar groups on the polymer and by a large organic surface area on the clay modifier. In fact, a large number of polymers, with increasing polarity, were tested, thermoplastic (PP, PS, SAN, PMMA, PVDF), and elastomeric (NBR, 33 wt.% AN), and the organoclay discussed in this section, having a large organic fraction, was reported to form intercalated nanocomposites (increase of d spacing) with any polymer but the apolar PP and PS. A primary alkyl ammonium, 12-aminolauric, was on the contrary not able to promote the intercalation of polymer chains in the interlayer space. Very similar XRD patterns were reported for cured compounds based on HNBR and Mt modified with octadecylamine, octadecyltrimethyl ammonium and with the polar ammonium cation discussed in this paragraph (methyltallow-bis-(2-hydroxyethyl)). Three detectable peaks are present in the pattern, the one at lower 2u value commented with polymer intercalation. The absence of peaks in the case of the uncured composite was commented with the exfoliated nature of the organoclay. 7.8 RUBBER NANOCOMPOSITES WITH MONTMORILLONITE MODIFIED WITH AN AMMONIUM CATION CONTAINING TWO LONG-CHAIN ALKENYL SUBSTITUENTS As summarized in Table 7.22, composites were prepared with Mt modified with a quaternary ammonium cation bearing two long-chain substituents, in most TABLE 7.22 Rubber Nanocomposites Based on Montmorillonite Modified with an Ammonium Cation Containing Two Methyls and Two Long-Chain Alkenyl Substituents Type of Clay

Type of Rubber

Blending Method

Na-Mt

NR

Melt Latex Solution Solution Melt Solution Melt Solution Melt Solution Melt

ENR BR SBR (H)NBR EPDM

References [16,17,28,29,38,40,42] [41] [20,30] [30] [17,47–49,53,55,56] [57] [17,56,82] [58,74] [92–97] [99] [104,105,114]

220

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

works either two tallow groups (preferentially hydrogenated, 2HT) or two stearyl groups. 7.8.1 In a Nutshell Data of the commercial organoclays employed for composites preparation are in Tables 7.23–7.28. From the recollection proposed in the following sections, some common features arise (i) A homogeneous dispersion of the organoclay in the rubber matrix was achieved, with exfoliated layers and stacks of few (even a couple) of them appearing in TEM micrographs. (ii) Two mechanisms were proposed for organoclay exfoliation: progressive delamination of clay stacks or chain intercalation followed by layers separation. (iii) In spite of the nice clay dispersion, XRD composite patterns showed in most cases (00l) reflections with some prevailing features, summarized below in points (iv)–(viii). (iv) In most cases several (00l) well-defined reflections were present in XRD pattern and, in some cases, the pattern of the composite was better resolved than the one of the starting organoclay. (v) In most cases, a shift of the (00l) peaks toward lower 2u angles was observed, passing from the organoclay to the composite. However, when an organoclay with a higher d001 (Dellite 67G) was used, the same interplane distances as in the organoclay were determined. (vi) When a shift was present, a larger interlayer distance was observed in composites based on a polar rubber. (vii) The peaks position did not depend on the organoclay content in the composite. (viii) Largely prevailing interpretation of experimental findings refer to polymer chains intercalation but an alternative view, that considers the intercalation of low molecular mass substances as likely to occur, was also proposed. (ix) Formation of the organoclay in situ, in the rubber medium was presented. 7.8.2 Composites with Two Talloyl Groups as Ammonium Cation Substituents 7.8.2.1 Comparison among Different Diene Rubbers Composites containing only the rubber and the organoclay (Dellite 67G from Laviosa) were prepared with many different rubbers [17,29]. As shown in Table 7.23, up to three well-defined and pretty narrow reflections were observed in composites patterns. The minor increase (if any) of d001 spacing from organoclay (3.6 nm) to composites (3.6–4.1 nm), the shape of the peaks, and the ratios between interplane distances

RUBBER NANOCOMPOSITES WITH MONTMORILLONITE MODIFIED

221

TABLE 7.23 RCN Based on Different Diene Rubbers and Mt Modified with an Ammonium Cation Containing Two Long-Chain Alkenyl Substituents (00‘) Interplane Distances in the Compositea Organoclay, d001 (nm)

Blending Method

Rubber

Dellite 67G, 3.6

–

Melt

1

2

3

4

References

–

3.6

1.84

1.22

0.89

[17,29]

NR IR BR E-SBRa S-SBRb S-SBRc

3.8 3.8 3.5 – – 3.9

1.92 1.96 1.8 2.06 2.0 1.9

1.3 1.26 1.2 1.4 1.3 –

– – – 1.04 1.0 –

a

Polybutadiene-co-styrene from emulsion polymerization. Solution polymerization with extension oil; trade name: SBR SOL R 72612. c Solution polymerization without extension oil; trade name: Buna 5025-0. b

led the authors to conclude that the ordered structures were due to a regular placement of alkenyl chains in the interlayer space and that polymer chains were not intercalated. Up to four well-defined reflections were detected in composites prepared by reacting pristine Mt and the ammonium chloride (with either saturated or hydrogenated tallows) [29,42], in situ in either NR or IR as the reaction medium, as shown in Table 7.24 and in Figure 7.9. They were analogously interpreted as due to the alkenyl chains arrangement, and the expansion of the d001 spacing (measured also as TEM micrographs), was commented with the formation of a perpendicular bilayer due to the intercalation of low molecular mass polar ingredients of natural rubber (see Section 7.9.4 for a more complete discussion). 7.8.2.2 NR as the Rubber In Table 7.25, data are summarized from XRD investigations on NR-based composites containing isolated lamellae and stacks of few or some of them. TABLE 7.24 RCN Prepared by Reacting, in NR or IR Rubber, Pristine Mt, and the Ammonium Chloride Containing Two Long-Chain Alkenyl Substituents (Either Saturated or Hydrogenated Tallows) (00‘) Interplane Distances in the Composites Blending Organoclay Rubber Method 1 (001) 2 (002) 3 (003) 4 (004) 5 (005) d001 References Mt þ 2HT NR IR Mt þ 2T NR IR

Melt

n.d. 4.0 n.d. n.d.

3.05 1.96 2.94 2.15

1.96 1.3 2.05 1.45

1.5 1.0 1.5 1.0

1.2 – 1.2 –

6.0 4.0 6.0 4.3

[29,42]

222

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

I (a.u.)

(001) 4.1 nm

(002) 2.0 nm (003) 1.3 nm (004) 1.0 nm (002) 3.0 nm (005) (003) (004) 2.0 nm 1.5 nm 1.2 nm

b

a 5

10

15 2θ (º)

20

25

30

FIGURE 7.9 XRD pattern of composites prepared by reacting pristine Mt and dimethyl-di (hydrogenated tallow) chloride, in situ in either NR (a) or IR (b) (see Refs [29,42]).

On the basis of TEM findings, it was hypothesized that the mixing process could induce only a partial delamination of the organoclay [38], whereas XRD pattern showed peaks in a higher number and better defined than in the organoclay, with 2u values independent on the clay content (from 5 to 20 wt.%). The same type of organoclay but with a lower interlayer spacing (about 3.1 nm, Cloisite 15A) was used in Refs [28] and [16], observed by TEM [16] a large prevailing clay exfoliation, with the absence of XRD reflections at low 2u angles [16,28], commented with polymer chains intercalation followed by clay exfoliation. Composites were prepared from a toluene solution of the organoclay and of a NR masterbatch with curatives [20]. A larger increase of the interlayer distance was found when a talloyl group was used in place of an ethyl–hexyl one (see Section 7.6.3) and was interpreted with the intercalation of polymer chains. Analogous findings and comments had been reported in Ref. [30], larger expansion with 2HT, due to polymer chains intercalation. 7.8.2.3 ENR as the Rubber The higher ability of a polar rubber to intercalate the OM–clay was invoked to comment a higher increase of d001 spacing, in the case of ENR 50-based composites [30]. 7.8.2.4 BR as the Rubber BR–organoclay composites prepared via melt blending showed the best defined XRD patterns among the ones based on an ammonium cation with two long-chain substituents: well-defined peaks [47–49,56], clearly detectable reflections [53], better definition, and a higher number of narrower peaks with respect to the starting organoclay [55,56]. Data are in Table 7.26. As for any other type of clay modifier, the pattern definition improved for samples processed on a two-roll mill or vulcanized (i.e., for samples containing

223

b

a

1.86 1.8–1.95

3.0, 1.1 3.15 3.196

d001 (nm)

Solution

Melt

Blending Method

Cloisite 15A from Southern Clay. From 2u values evaluated from the published XRD pattern.

– –

125 125a 125a

a

Mt CEC (mequiv./100 g)

Organoclay

ZnO, SA, S, CBS, ZnO, SA, S, CBS, Ca stearate, MBT, CBS

ZnO, SA, S, MBTS, PBN ZnO, SA, S, MBTS, PBN ZnO, SA, S, CBS, NS

Further Ingredients

3.53 3.0–3.16

– – 4.4–4.9

1

– – 1.5b – –

1.65b –

3 – – 2.2–2.4b

2

(00‘) Interplane Distances in the Composites

[20] [30]

[16] [28] [41]

References

TABLE 7.25 RNC Based on NR and Montmorillonite Modified with an Ammonium Cation Containing Two Long-Chain Alkenyl Substituents

224

d

2.4 3.37 1.9, 1.3

1.89

3.1 2.4 3.5, 1.8, 1.2

Organoclay prepared in situ

3.32, 1.8b, 1.2b 2.4 3.56, 1.80, 1.20

d001 (nm)

Solution

Melt

Solution

Melt

Blending Method

1.8c 1.87 1.9c

3.2c 3.89 4.1

ZnO, SA, S, TMTD ZnO, SA, IPPD, S, CBS Curatives

SA, S, CBS unvulvanized SA, S, CBS vulcanized SA, IPPD, CZ, S

b

a

1.2c 1.21 1.3c

– – – 1.2c 1.3c 1.33c –

2.3c (002) 2.05 1.92 1.8c 1.9c 2.0c – 5.07 (001) 4.2 4.1 3.61 3.85 3.99 3.63

SA, S, CBS vulcanized SA, S, MBTS SA, S, MBTS

ZnO, ZnO, ZnO, – ZnO, ZnO, ZnO,

1.9c (002)

1.5 (003) 1.2c 1.36c (003), 1.0c (004) 1.3c (003)

3

2.2 (002) 1.83c 2.0c (002)

2

4.17 (001)

4.75 (001) 3.8c 4.20 (001)

1

(00‘) Interplane Distances in the Composites

ZnO, SA, S, CBS unvulcanized

DCUP ZnO, SA, S, TMTD ZnO, SA, S, CBS unvulcanized

Further Ingredients

Nanolin DK4 (40% 2Me2HT) from Zhejiang Fenghong Clay Chemicals Co., Ltd, China. Cloisite 20A from Southern Clay. c From 2u values evaluated from the published XRD pattern. d Cloisite 15A from Southern Clay.

119 SBR 95b 125 –

125 95b –

BR –a 95b –a

Mt CEC (mequiv./ 100 g)

Organoclay

[56] [82] [74]

[57]

[53]

[55]

[47] [56] [48,49]

References

TABLE 7.26 RNC Based on BR, SBR, and Montmorillonite Modified with an Ammonium Cation Containing Two Long-Chain Alkenyl Substituents

RUBBER NANOCOMPOSITES WITH MONTMORILLONITE MODIFIED

225

oriented clay stacks). Broad peaks were instead present in patterns of uncured masterbatches from an internal mixer [55] and of composites prepared from a rubber solution [57]. Expansion of interlayer distance from organoclay to composites was observed and shape and position of (00l) reflections in composites were such as to allow their interpretation, in some cases, as successive order basal reflection [47–49]. Peaks position was not found to change with the organoclay content (5–30 phr) [47]. The in situ preparation of the organoclay in the rubber matrix was proposed in Refs [48,49] and, analogously to what reported in Refs [17,29,42], the pattern of the composite obtained with the preformed organoclay was reproduced. 7.8.2.5 SBR as the Rubber Besides the excellent clay dispersion (as for the other rubbers), strain-induced orientation of the anisotropic particles was observed [58]. Very similar values of interplane distances were detected for the organoclay in SBR-based composites, prepared from either melt [56,82] or emulsion blending (see Table 7.26). Values and in some cases also shape of XRD patterns were very similar to those observed in the BR matrix. As reported above, an intercalated structure with only low molecular mass substances in the interlayer space was hypothesized in Ref. [17], on the basis of TEM and XRD findings very similar to those proposed in the other works. 7.8.2.6 NBR as the Rubber With NBR as the rubber, most silicate layers were observed as monolayers, with some couple of layers joined together in the intercalated form [93]. Also for this rubber, composites XRD patterns showed in most cases three peaks and were even with a higher number and better defined peaks than the ones of the starting organoclays [94]. Data are shown in Table 7.27. The interlayer distance was found to increase to values lying in the same range reported for the other rubbers, being larger with a larger amount of chemicals in the compound, the lowest value was in the presence of only a peroxide [95]. Also the lower detected interplane distances were in a pretty narrow range: 1.84–1.97 and 1.26–1.34 nm. The interlayer distance was found to increase with NBR polarity, that was thus commented, as in the case of ENR, to favor a larger polymer intercalation [94]. The effect of vulcanization on interplane distance was investigated. A slight expansion [93,96] was commented with an advanced intercalation during the vulcanization process. A decrease of d001 spacing [94,97] was attributed to deintercalation and to chain expulsion due to vulcanization pressure. Moreover, the intercalation of polymer chains was hypothesized to occur thanks to the interaction between the ammonium cation and the crosslinking precursor attached to the polymer [110]. The large organic surface area provided by 2HT was considered as responsible for the polymer intercalation [92]. 7.8.2.7 EPDM as the Rubber Silicate layers were uniformly dispersed in EP (D)M as monolayers or a few layers, but clay agglomerates were reported to exist in composites with MAH-g-EPDM [104]. Data from XRD analysis are collected in Table 7.28. XRD composite pattern was better defined (with more peaks) than the one obtained with a trimethylammonium cation [104]; expansion of interlayer distance was attributed to the presence of C18 substituents, favoring the polymer intercalation.

226

–

2.42 2.8

125a

95c 93d modifier content 93d modifier content

3.15, 1.26b

125a

Solution

Melt

Blending Method Further Ingredients

b

3.53–4.26

4.12 4.35 4.70 3.69 4.0, NBR, 42% ACN 3.7, NBR, 42% ACN

ZnO, SA (1 phr), S, ZDDC ZnO, SA (2 phr), S, ZDDC ZnO, SA (4 phr), S, ZDDC DCUP ZnO, SA, S, MBTS, uncured ZnO, SA, S, MBTS, cured ZnO, SA, S, accelerators

3.57

4.2b 3.2b 3.9b 3.6

1

1.76–2.1b

1.90 1.97 – 1.88 2.0b 1.8b

1.84

2.0b 1.6b 1.9b –

2

1.2–1.4b

1.29 1.34 – 1.26 1.3b –

1.26

1.4b 1.1b – –

3

(00‘) Interplane Distances in the Composites

–

ZnO, SA, S, TBBS, unvulcanized ZnO, SA, S, TBBS, vulcanized – DCUP

Cloisite 15A from Southern Clay. From 2u values evaluated from the published XRD pattern. c Cloisite 20A from Southern Clay. d Nanofil 15 from Southern Clay. e Nanomer I44P from Nanocor.

a

2.2

–e

2.98, 1.26

d001 (nm)

Mt CEC (mequiv./ 100 g)

Organoclay

[99]

[94]

[93,96]

[92] [95]

[97]

References

TABLE 7.27 RNC Based on NBR and Montmorillonite Modified with an Ammonium Cation Containing Two Long-Chain Alkenyl Substituents

227

, 95 , 120

2.4 3.4, 1.9, 1.27

4.2, 2.1, 1.4

d00‘ (nm)

Melt

Blending Method

– ZnO, SA, MBT, CBS, S

ZnO, SA, S, CBS

Further Ingredients

b

Cloisite 20A from Southern Clay. From 2u values evaluated from the published XRD pattern. c Dellite 67G from Laviosa.

a

a

a

110

Mt CEC (mequiv./ 100 g)

Organoclay

5.5 4.1, EPDM/MAH-g-EPDM 2.7 3.4 4.4a, EPDM/MAH-g-EPDM

1

2.2 2.05 1.3a 1.9 1.9a

2

1.4 1.36 – 1.27 1.27a

3

(00‘) Interplane Distances in the Composites

[114] [113]

[104]

References

TABLE 7.28 RNC Based on EPDM and Montmorillonite Modified with an Ammonium Cation Containing Two Long-Chain Alkenyl Substituents

228

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

On the contrary, in Ref. [113] a larger d spacing was reported only in the presence of EPDM-g-MAH and it was concluded that the intercalation of polymer chains could occur only thanks to the grafted compatibilizer. Also in the case of a TPV (an EPDM/ PP blend) [114], the interlayer space was found to change to a very minor extent, from 2.4 to 3.7 nm.

7.9 PROPOSED MECHANISMS FOR THE FORMATION OF RUBBER–CLAY NANOCOMPOSITES RCN are thus characterized by the presence of (i) clay stacks with organic substances intercalated between two opposite layers and a variable number of layers in a stack; as reported in Section 7.2.4, they are traditionally named as intercalated organoclay; (ii) isolated clay lamellae: clay is named in this case as exfoliated or delaminated. Two mechanisms are available in the open literature for explaining the formation of both (i) and (ii). 7.9.1 Two Mechanisms for the Formation of an Exfoliated Clay Two mechanisms are available in the literature to justify the formation of an exfoliated clay: (i) Intercalation of polymer chain in the interlayer space with subsequent clay exfoliation. (ii) Progressive peeling off of clay stacks through a so-called “delamination mechanism.” 7.9.2 Two Mechanisms for the Formation of an Intercalated Organoclay Two mechanisms are available in the literature to justify the formation of an intercalated organoclay: (i) Polymer chains are intercalated in the clay interlayer space, with the expansion of the interlayer distance and, in most cases, the formation of a crystalline intercalated organoclay with reflections detectable in the XRD pattern. (ii) The intercalation of polymer chains is considered as highly unlikely and the organic moiety between two opposite layers is supposed to be made by low molecular mass substances: ammonium cations and chemicals. The variation of the interlayer distance is supposed to be due to reorganization of

PROPOSED MECHANISMS FOR THE FORMATION OF RUBBER–CLAY NANOCOMPOSITES

229

the ammonium cations substituents as well as to the other cointercalated low molecular mass substances. These mechanisms are examined in some detail in the following sections. Mechanism (ii) evidently implies clay exfoliation to occur through a progressive layer separation from a clay stack. 7.9.3 Intercalation of Polymer Chains in the Interlayer Space Figure 7.10 summarizes mechanism for polymer intercalation, with reference to an XRD pattern: a (00l) reflection shifts toward lower 2u angles as the interlayer space increases, as a consequence of polymer chains intercalation. As mentioned above, most literature refers to this mechanism to explain any shift of the (00l) peaks toward lower 2u angles. A summary can be attempted as follows (i) There is a general agreement on the chance for polymer chains to enter in the interlayer space of a clay modified with either an amine or an ammonium salt. (ii) In the case of a pristine clay, the major part of papers is prone to exclude this phenomenon.

I (a.u.)

As mentioned in Section 7.2.4, the elaborations on rubber–clay composites occurred after the publication of some fundamental papers on composites based on

4

5

6

7 2θ (º)

8

9

10

FIGURE 7.10 Schematic representation of progressive change that can occur to the (001) reflection of a clay as a consequence of mixing with polymer matrices: clay intercalation with polymer molecules eventually leading to clay exfoliation.

230

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

thermoplastic materials. In particular, the 1993 Giannelis paper [118] appears as a milestone in the field. In this paper, and in the related patent [119] it is reported and claimed as follows: (i) To have the intercalation of an apolar polymer such as polystyrene into the clay interlayer space, the clay has to be organically modified. (ii) The polystyrene composite with an Mt modified with a dimethyl dioctadecyl ammonium bromide that showed peaks shifted toward lower 2u (d001 from 2.52 to 3.2 nm). This milestone paper and the subsequent first reviews [1–3] paved the way for the interpretation of RCN morphology, summarized in this section. As far as the clay modifier is concerned, its organic surface is reported to have an effect in favoring the intercalation of polymer chains, as well as its reactivity with curatives. A characteristic step of rubber technology, such as curing, is also reported to affect the interlayer distance, in most cases reducing it through the deintercalation of modifiercurative(s) reaction product(s). Pressure applied in the curing step, favoring clay orientation, appears to be beneficial for a better XRD pattern definition. In many reports, comments on polymer chains intercalation are also supported by TEM analysis, by observing the presence of an organic moiety between two opposite layers and also by comparing the interlayer distances, measured through XRD and TEM micrographs. Authors in favor of this mechanism hypothesize that clay exfoliation essentially occurs through polymer intercalation. 7.9.4 Intercalation of Low Molecular Mass Substances in the Interlayer Space The intercalation of essentially low molecular mass substances was proposed [29] on the basis of experimental findings, essentially from XRD and TEM analysis, from RCN obtained (i) by blending rubber and an organoclay; (ii) by preparing in situ the organoclay in the rubber medium, by reacting a pristine clay and the ammonium cation; (iii) by preparing in situ organoclays with further low molecular mass chemicals as guests in the interlayer space. The above-mentioned points are discussed in more detail in the following paragraph. Data reported in Table 7.23 showed interlayer distances considered essentially the same in the starting organoclay (Dellite 67G) and in the composites based on many different rubbers [9–19,45,47–51,59–61,102]. It was also shown in Refs [29,42,43] that the same interlayer distance (about 4 nm) was obtained by preparing the organoclay in situ, in many different hydrocarbon polymers, through the reaction of pristine Mt with an ammonium cation (see Table 7.29). By performing the in situ modification of pristine Mt in NR [29,42], a larger interlayer distance (about 6 nm

231

Rubber

IR PE EPDM BIIR NBR

NR IR PE EPDM BIIR NBR

Organoclay

Mt þ 2HT

Mt þ 2HT þ SA

Melt

Melt

Blending Method

n.d. n.d. n.d. n.d. n.d. n.d.

4.0 4.17 – – –

1 (001)

3.05 3.0 – 3.0 3.0 3.0

1.96 2.0 1.99 1.96 1.91

2 (002)

1.96 2.0 2.0 2.0 2.0 2.0

1.3 1.3 1.31 1.32 1.27

3 (003)

1.5 1.5 1.5 1.5 1.5 1.5

1.0 1.0 1.0 0.98 0.98

4 (004)

1.2 1.2 1.2 1.2 1.2 1.2

5 (005)

(00‘) Interplane Distances in the Composites

– 1.0 – 1.0 – –

6 (006)

6.0

3.9

4.0

d001

[29,42] [43]

[29,42] [43]

References

TABLE 7.29 RCN Prepared by Reacting Pristine Mt and the Ammonium Chloride Containing Two Long-Chain Alkenyl Substituents (2HT) in the Presence or in the Absence of Stearic Acid (SA)

232

MORPHOLOGY OF RUBBER–CLAY NANOCOMPOSITES

3.0 nm

2.0 nm 1.5 nm 1.2 nm

I (u.a.)

a

b

c d

0

5

10

15

20

25

30

2θ (º)

FIGURE 7.11 X-ray diffraction (Cu Ka) pattern, in the 2u range 2–30 , of the composite obtained by mixing four different polymers: (a) PE; (b) EPDM; (c) BIIR; (d) NBR with pristine Mt (5.7 php), the ammonium salt (dimethyl-di(hydrogenated tallow chloride) (4.2 php), and stearic acid (6.0 php) (see Ref. [43]).

instead of about 4 nm) was detected; authors speculated on the presence in the interlayer space of low molecular mass chemicals present in NR as cointercalated guests. In Ref. [43], interplane distances were compared of organoclays obtained in situ by reacting pristine Mt and 2HT as the ammonium cation, in many different hydrocarbons, in the presence or in the absence of stearic acid, a striking reproducibility of interplane distances is shown in Table 7.29 and in Figure 7.11. Author proposed that the organoclays with basal spacing of about 4.0 nm possibly corresponded to paraffin-type tilted bilayer intercalates, whereas organoclays with basal spacing of about 6.00 nm were due to paraffin-type perpendicular bilayer intercalates, with stearic acid as a guest It was commented that the large cross-sectional area of the double-chain amphiphile was further increased by the parallel placement of the long alkyl chain of stearic acid, thus leading to their common orientation perpendicular to the clay layers. This mechanism suggests thus to consider the expansion of a clay d basal spacing as not necessarily due to polymer chains intercalation. In the frame of this hypothesis, the presence of nicely resolved peaks, with a better definition of composite patterns with respect to the organoclay ones, the periodicity of 2u angle values, are considered as meaningful supporting indications. ABBREVIATIONS CB CBS

carbon black N-cyclohexyl-2-benzothiazole sulfenamide

REFERENCES

CDDC DCUP IPPD MBT MBTS PBN 6PPD S SA sDPA TAIC TBBS TMTD TOTM ZDC ZDDC ZDEC ZMMBI

233

copper dimethyl dithio carbamate dicumyl peroxide N-isopropyl-N0 -phenyl-p-phenylene diamine 2-mercaptobenzothiazole 2,20 -dithio dibenzothiazole phenyl-b-naphthylamine N-(1,3-dimethylbutyl)-N0 -phenyl-p-phenylene diamine sulfur stearic acid substituted diphenyl amine triallyl isocianurate N-tert-butyl-2-benzothiazole sulfenamide bis(dimethylthiocarbamyl) disulfide trioctyl trimellitate 77, 78 zinc dithio carbamate zinc dimethyl dithio carbamate zinc diethyl carbamate zincmethylmercaptobenzimidazole

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CHAPTER 8

RHEOLOGY OF RUBBER–CLAY NANOCOMPOSITES AVRAAM I. ISAYEV JAESUN CHOI

8.1 INTRODUCTION Polymer nanocomposite fabrications have been considered as one of the novel technologies leading to manufacturing of novel products. Numerous studies have been carried out to develop nanocomposites with an enormous enhancement of material properties. In the rubber industry, the term “nanocomposite” does not sound innovative, since the conventional reinforcing particles being used in the rubber industry, such as carbon black and silica, are of a nanosize dimension depending on the particle grade. Therefore, it is not surprising the conventional rubber composites containing carbon black and/or silica being used for many years are the original “nanocomposites.” However, the research and development of new nanoparticles such as nanoclays have been carried out during last few decades. The outstanding features of nanoclays are their platelet shape with a high aspect ratio and a nanosize thickness. These platelets are stacked with other platelets forming agglomerate. However, if stacked platelets can be exfoliated such that stacks are separated to individual platelets, interacting strongly with rubber matrix, one can obtain the nanocomposites with significantly enhanced properties. Then, it will be possible to reduce the concentration of nanoparticles in composites such that the target properties can be met at a lower amount of nanoparticles, providing tremendous benefits compared to the conventional composites. In particular, with relevance to future light-weight tires that will be made of rubber nanocomposites, a less consumption of energy is expected. The rheological properties of rubber nanocomposites are important in manufacturing of various products. In particular, the processing, morphology and mechanical properties of rubber–clay nanocomposites are defined by the rheological properties. A significant increase in viscosity is expected due to the addition of nanoclays into rubber Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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matrix. This is caused by a reinforcing effect and a rubber–clay interaction through the intercalation and exfoliation of nanoclays. The degree of exfoliation of nanoclay can be related qualitatively to the extent of the increase in viscosity. A review paper [1] devoted to the melt rheology of polymer–clay nanocomposite is very helpful in understanding the rheological behavior of polymer nanocomposite. Also, a recent book [2] specifically provides the summary of current understanding on rheology of rubber nanocomposites. There are other review papers devoted to rubber–clay nanocomposites providing a broad understanding of their properties [3–5]. Although numerous research groups have been involved in the investigation of rubber–clay nanocomposites, the study devoted to rheology are surprisingly lacking. In this chapter, studies related to the rheological behavior of rubber–clay nanocomposites were reviewed with aim to provide better understanding of the subject matter. There are number of papers dealing with rheological behavior of various rubber–clay nanocomposites and providing some understanding of rheological behavior [6–15]. The studies devoted to the rheological behavior of the clay composites containing different rubbers are described including natural rubber (NR) [16], polyisoprene rubber (IR) [6,17,18], epoxidized natural rubber (ENR) [18], styrene–butadiene rubber (SBR) [7,19], polybutadiene rubber (BR) [7,8,20], acrylonitrile butadiene rubber (NBR) [7,21,22], ethylene propylene diene monomer (EPDM) [9,23–26], fluoroelastomer [10,27], poly(isobutyleneco-para-methylstyrene) (BIMS) [28], poly(ethylene-co-vinylacetate) (EVA) [11,29], polyepichlorohydrin (PECH) [12], thermoplastic polyurethance (TPU) [13,30,31], SBS, or styrene–ethylene–butylene–styrene (SEBS) block copolymer [14,32,33], and rubber/plastic blends [15,23,34–37].

8.2 RHEOLOGICAL BEHAVIOR OF RUBBER–CLAY NANOCOMPOSITES In general, the addition of a filler into a rubber matrix causes an increase in viscosity of the composite. This behavior is caused by a hydrodynamic effect due to a volume fraction of the filler in a composite, a physicochemical rubber–filler interaction, and a filler–filler interaction known as Payne effect. The most papers summarized in this section reported the increase in viscosity or the storage modulus showing pseudosolid-like behavior by adding clay in rubber matrix. In addition, several rheological characteristics such as the power-law index, consistency index, activation energy, and percolation threshold were obtained from the analysis of rheological behaviors. The observed increase in viscosity or the storage modulus was explained by the improved rubber–clay interaction, due to the clay intercalation and exfoliation. The rheological behavior of rubber–clay nanocomposites is more complex than that of the conventional rubber composites containing carbon black and silica, since the orientation of clay layers at high shear rates may lead to the reduction of viscosity. Therefore, a competition between clay filling and orientation would dictate rheological behavior of rubber–clay nanocomposites. The following sections provide specificity of rheological behavior of nanocomposites based on type of rubbers.

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8.2.1 Natural Rubber (NR), Epoxidized Natural Rubber (ENR) and Polyisoprene Rubber (IR)–Clay Nanocomposites Vu et al. [18] studied structures and properties of synthetic cis-1,4-polyisoprene rubber and ENR composites filled with clays at the loadings of 10, 20, and 30 phr. The ENR contained epoxy groups of 25 mol% (ENR25) and 50 mol% (ENR50). Several clays were used in this study. They were the products of Southern Clay, including pristine montmorillonite and three organoclays modified with bis(2-hydroxyethyl) methyl tallow amine (C1), dimethyl dehydrogenated tallow amine (C2), and dimethyl hydrogenated tallow 2-ethylhexyl amine (C3). Melt mixing was carried out in an internal mixer for IR and ENR25 composites, while a solution mixing was used for ENR50 composites. Rheological behavior in strain sweep mode was obtained using RPA 2000 at a temperature of 50
C and frequency of 5 Hz. For the ENR50–clay composites at the loading of 30 phr, the composite containing the clays C1 and C2 showed a strong Payne effect and a higher storage modulus at low strains, while the composite containing the pristine clay showed a very little Payne effect and a low storage modulus at low strains. These behaviors were ascribed to the existence of a strong interaction between polar rubber matrix (ENR50) and organoclay (C1, C2), leading to the better intercalation as confirmed by XRD results. At the same time, the increased amount of bound and occluded rubber acted as an additional filler increasing the effective volume fraction of filler and subsequently generating the stronger clay networks, as shown in the extent of Payne effect. The storage modulus at low strains of composites containing the clay C2 at the loading of 10 phr in the polar matrix of ENR50 was higher than that of silica (Hi-Sil233, Harwick Co.) filled composite at the same loading. However, the composite of the nonpolar matrix IR with an addition of 10 phr of silica showed a higher storage modulus than that with clay C2. This observation was in contrast to ENR50 composite, indicating the important role of polarity of rubber matrix to the intercalation and rubber–clay interaction. Jeon et al. [6,17] studied rheological behavior of IR–clay nanocomposites. IR was solution blended with Na-montmorillonite (Cloisite Naþ , Southern Clay) and organophilic montmorillonite modified with methyl tallow bis-2-hydroxyethyl ammonium (Cloisite 10A, Southern Clay). The rheological properties in frequency sweep were obtained using an oscillatory rheometer with parallel-plates geometry (Physica UDS 200) at a temperature of 50
C under N2 flow and a strain amplitude below 2%. The storage modulus, G0 , of the IR–organoclay nanocomposite containing 5 wt.% of clay at a low frequency (v ¼ 0.1 rad/s) was six orders of magnitude higher than that of unfilled IR, while the storage modulus of the IR–pristine clay nanocomposite was just an order of magnitude higher. The huge increase in the storage modulus for the IR–organoclay composite was ascribed to the high degree of clay exfoliation, as confirmed by XRD as well as SAXS studies, and the network formation of clay platelets. The plots of log G0 versus log v presented in Figure 8.1a indicate the increase in the storage modulus in the all range of frequency with the loading of organoclay. The slopes of G0 versus v curves in the low-frequency range (v < 0.5 rad/s) decreased gradually with the loading of organoclay up to 4 wt.%.

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106 5

10

105 3

η ′ (Pa. s)

G′ (Pa)

10

101 10–1 10–3

104 103 102

10–1

100

101

ω (rad/s)

102

101

10–1

100

101

102

ω (rad/s)

FIGURE 8.1 Storage modulus (a) and dynamic viscosity (b) versus frequency for IR–organoclay composites at 50
C and various contents of organoclay in wt.%: 1 (circle), 2 (square), 3 (diamond), 4 (cross), 5 (plus), 7 (triangle), 9 (reverse triangle) [6].

However, the slopes were similar above the loading of organoclay of 4 wt.%, indicating a transition from the liquid-like to pseudo solid-like behavior of nanocomposites. The clay loading about 4 wt.% corresponding to the volume fraction of 1.9%, was considered as the concentration where the rheological percolation in terms of the storage modulus takes place in IR–clay nanocomposite. A Newtonian behavior of dynamic viscosity is seen at the organoclay loadings of 1 and 2 wt.% (Figure 8.1b). As the loading of organoclay increased, the shear thinning behavior instead of the Newtonian behavior, was observed at any frequency. Although the dynamic viscosity does not strictly follow the power-law behavior, authors determined the power-law index (n) by fitting a power-law Eq. (8.1) to the viscosity curve. h ¼ K g_ n  1

ð8:1Þ

where, h is the viscosity, K is the consistency index, g_ is the shear rate, and n is the power-law index. The power-law index decreased with loading of clay and reached saturation at a loading of 5 wt.%. The values of the power-law index reported were 0.65, 0.34, and 0.17 at 3, 4, and 5–9 wt.%, respectively. Viscosity percolation threshold was obtained from the plots of relative dynamic viscosity hr (hr ¼ h0 /h0 at v ¼ 0.1 rad/s, h0 is zero-shear viscosity of IR) versus filler volume concentration, showing the viscosity percolation threshold of about 7 wt.% corresponding to 3.2 vol% which is higher than that based on the storage modulus. To describe the relative dynamic viscosity as a function of the filler volume concentration, w, the modified Krieger’s empirical model equation [38] was applied:  hr ¼ 1 

w wm;e

 ½hwm;e

ð8:2Þ

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245

where wm,e is the effective maximum packing volume fraction, [h] is the intrinsic viscosity. The intrinsic viscosity was obtained by the fitting of this equation to the dependence of hr on (1  w/wm,e), assuming wm,e is equal to the volume concentration of percolation threshold. The calculated value of [h] was used to determine the aspect ratio of clay in Eq. (8.3) below, based on the Douglas and Garbochi approximation obtained from the Monte Carlo simulation of random dispersion of oblate ellipsoids [39]: ½ h 

1012 þ 2904b  1855b1:5 þ 1604b2 þ 80:44b3 1497b þ b2

ð8:3Þ

where b is the ratio of the length of the minor to major symmetry axis of the ellipsoid. The calculated aspect ratio of clay (1/b) was approximately 282, which was different from the result of TEM image analysis showing the aspect ratio of approximately 110. The aspect ratio of clay in composite was also calculated [17] based on a Pade-type approximation [39]. wp 

9:875b þ b2 7:742 þ 14:61b þ 12:33b3=2 þ 1:763b2 þ 1:658b3

ð8:4Þ

where, wp ¼ 0.02 is the volume concentration of percolation threshold based on the storage modulus, b is the inverse of the aspect ratio of the oblate ellipsoids. The calculated average aspect ratio of clay in IR composite was 62 by this approach, which was lower than the measured value by TEM image analysis. The interesting observation was made based on a dependency of the storage modulus of IR–organoclay composites on preshear [17]. The IR–organoclay composites at the loading of 3 wt.% (below percolation) and 5 wt.% (above percolation) were subjected to preshear at a shear rate of 1 s1 before starting the rheological test. As shown in Figure 8.2, the storage modulus of the presheared composite containing 5 wt.% that is above percolation decreased in the whole frequency range compared to that without preshear. However, the storage modulus of the presheared composite containing 3 wt.% that is below percolation was significantly higher than that without preshear. In the case of the composite containing organoclay above the percolation threshold, the decrease in the storage modulus by preshear was caused by deformation of the clay network structures due to the shear-aligned clay platelets. The large increase of the storage modulus by preshear at 3 wt.%, which is below the percolation threshold, was probably due to the better shear-induced interaction and dispersion. Stephen et al. [16] studied rheological behavior of NR latex–clay nanocomposites. NR latex was mixed with sodium bentonite (EXM 757, Sud Chemie) and sodium fluorohectorite (Somasif ME-100, Coop Chemicals). The concentration of clays varied from 1 to 2.5 phr. The rheological properties were obtained using a coaxial cylinder viscometer (Haake Viscotester VT 550) at a temperature of 25
C. All the composites exhibited the increase in viscosity with the loading of clays. This

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FIGURE 8.2 Effect of shear-induced alignment on the storage modulus of the IR–organoclay composites at 50
C and loadings of 3 wt.% (circle) and 5 wt.% (triangle). The open and filled symbols correspond to the unaligned and aligned samples, respectively [17].

behavior corresponded to the XRD data of the composites, showing the intercalation of rubber into the clays. In order to evaluate the rubber–clay interaction the Kraus plot [40] was used. This is the plot of vr0/vr versus w/(1  w), where vr0 is the volume fraction of rubber in the swollen gel of gum vulcanizate, vr is the volume fraction of rubber in the swollen gel of filled vulcanizate, w is the volume fraction of clay of the composite. The slopes of the Kraus plots were negative with the value of  8.96 and  2.09 for the bentonite- and fluorohectorite-filled nanocomposite, respectively, at the loading of 2.5 phr. These negative values of slopes indicate the rubber–clay interaction, which was stronger in the bentonite-filled composite. The power-law index was calculated by fitting the viscosity curve to the power-law equation. The power-law index was 0.23, 0.87, and 0.92 for the 2.5 phr of bentonite, fluorohectorite-filled composites at the loading of 2.5 phr and unfilled NR latex, respectively. The pseudo solid-like behavior was observed in the bentonite-filled composites at the loading of 2.5 phr, while the unfilled NR latex showed the behavior close to that of the Newtonian fluid. The higher viscosity and the lower power-law index of the bentonite-filled composites were due to the stronger rubber–clay interaction as supported by the slope of the Kraus plot and the interlayer spacing from XRD data. 8.2.2 Styrene–Butadiene Rubber (SBR)–Clay Nanocomposites Tian et al. [19] studied structures and properties of SBR–fibrillar silicate composites. The material used in their study was bis-(3-triethoxysilylpropyl)-tetrasulfide (TESPT) and natural fibrillar silicate provided by Jiangsu AT Co. Rheological behavior in strain sweep mode was obtained using RPA 2000 at a temperature of

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60
C and a frequency of 1 Hz. SBR was compounded with clay at a loading of 40 phr in a two-roll mill. The drop of the storage modulus with strain up to 400% for the SBR–clay composite containing TESPT was significantly lower than that of the composite without TEPST. It was ascribed to the presence of a relatively weak filler network in the SBR–clay–TESPT composite, due to the uniform dispersion of clay which is in agreement with the results of morphology tests. The SBR–clay composite after further mixing showed the less storage modulus drop with strain compared to that of the composite without further mixing. This behavior indicated the loss of the clay network due to the breakage of the fibrillar silicate upon the further duration of mixing. Sadhu and Bhowmick [7] studied rheological behavior of SBR–clay composites and reported very unusual results. The organically modified montmorillonite was prepared by treating Na-montmorillonite (Cloisite Naþ , Southern Clay) with octadecylammonium chloride (C18). The organoclay was mixed with SBR by a solution mixing at loadings of 2, 4, 6, and 8 phr. Viscosity of the composites was measured using a capillary rheometer (Monsanto Processability Tester) at a temperature of 130
C. The significant viscosity reduction of the nanocomposites was observed with the addition of organoclay up to 8 phr. In fact, viscosity of all the SBR–organoclay composites was lower than that of unfilled SBR. According to the authors, the decrease in viscosity by adding the organoclay was due to the better orientation of the clay platelets under the high shear and the slippage of rubber chains on the clay surface. However, viscosity of the SBR–pristine clay composite at the loading of 4 phr was similar to that of unfilled SBR. The extrudates of SBR–organoclay exhibited the more smooth surface and lower die swell than those of SBR–pristine clay composite at the loading of 4 phr. Stephen et al. [16] studied rheological behavior of carboxylated SBR–latex nanocomposite as well as NR–latex nanocomposites discussed earlier. Viscosity increased with the loading of clays up to 2.5 phr for the bentonite- and fluorohectorite-filled composites. However, viscosity of the fluorohectorite-filled composites was lower than that of unfilled SBR latex at a high shear rate of 190 s1, while it was higher at low shear rates, below 90 s1. The calculated power-law index was 0.32, 0.44, and 0.72 for the bentonite-, fluorohectorite-filled composites at the loading of 2.5 phr and unfilled SBR latex, respectively. The slope in the Kraus plots [40] was 19.5 and 7.0 for bentonite- and fluorohectorite-filled composites, respectively, at a loading of 2.5 phr. Accordingly, the SBR latex–bentonite composite exhibited a stronger rubber–clay interaction than the fluorohectorite-filled composite, leading to the higher viscosity and more tendency to the pseudo solid-like behavior. 8.2.3 Polybutadiene Rubber (BR)–Clay Nanocomposites Sadhu and Bhowmick [7] studied rheological behavior of BR–clay composites. The composite preparation and experimental methods were the same as summarized earlier in case of SBR–clay composite. The rheological behavior of BR/clay composites was quite different from that of SBR–clay composites [7]. When organoclay at a loading of 4 phr was added to BR matrix, viscosity of the composite

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slightly increased. However, by adding pristine clay at the same loading, viscosity of the composite was similar to that of unfilled BR. It is not clear why the BR–pristine clay composite did not show the increase in viscosity, although the XRD data [41] revealed the exfoliation of the pristine clay in the composites. The extent of die swell was lower for the BR/pristine clay compared to the BR–organoclay composite at a loading of 4 phr. Wang et al. [8] studied rheological behavior of liquid BR–clay composites and reported an unusual behavior. The BRs used in their study were 1,2-polybutadiene oligomer without hydroxyl end groups (PB) and hydroxyl-terminated 1,4-polybutadiene oligomer (HTPB) with Mn of 2200 and 4200. Organoclay was prepared by treating pristine montmorillonite, supplied by Tianjin Organic Clay Co., with octadecyltrimethylammonium chloride (C18). The composites were prepared by the solution mixing with clay at a concentration of 5 wt.%. The rheological properties were obtained using the rheometer with parallel plates (HAAKE Rheo-Stress 600) in both steady and oscillatory flows at a temperature range of 26–136
C and a heating rate of 2
C/min and a shear rate of 1 s1 under N2 flow. The specimens were subjected to the shear rate of 1 s1 for a period of 15 min before heating. Here, the results of HTPB–clay composite with matrix of higher Mn are summarized, since the behavior of composite with matrix of lower Mn was similar. As shown in Figure 8.3 for the HTPB–clay composite, the decrease in viscosity with temperature with viscosity minimum at 60
C is observed. After passing through the minimum point, viscosity increases up to 100
C. Then, viscosity remained a constant and slightly decreased as temperature increased to 136
C. However, this behavior was not observed in PB–clay composite with matrix not containing hydroxyl groups (inset in Figure 8.3).

FIGURE 8.3 Temperature dependence of the steady shear viscosity for HTPB4200–organoclay and HTPB2200–organoclay nanocomposites at a loading of 5 wt.%. The inset shows temperature dependence of steady shear viscosity for (a) virgin HTPB2200, (b) HTPB4200, (c) PB, (d) PB–organoclay (5 wt.%) composite [8].

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249

Therefore, the presence of hydroxyl end groups in HTPB was responsible for the increase in viscosity upon heating. The temperature where viscosity began to increase was called a transition temperature. The increase in viscosity after the transition temperature was due to the occurrence of the strong polymer–clay interaction between the polar groups and the subsequent exfoliation of clay. Above a temperature of 100
C, the temperature increase and the shear-induced orientation of the clay platelets, leading to a decrease in viscosity, contributed more to the reduction of viscosity than the interaction between polymer and clay, leading to an increase in viscosity. The saturation and slight decrease of viscosity at higher temperatures was explained by these competing effects. After temperature reached 136
C in the heating cycle, the sample was further subjected to shear for a period of 2 min, then temperature decreased to 26
C. The reduction of viscosity is observed during cooling from 136 to 80
C (Figure 8.3), due to the shear-induced orientation of the exfoliated clay leading to the decrease in viscosity. Viscosity at 26
C in the cooling process was significantly higher than that initially measured. This difference in viscosity was due to the clay exfoliation, which is irreversible in heating and cooling cycles. The effect of shear rate on viscosity was also evaluated for HTPB–clay composite containing 5% of organoclay. Viscosity was measured at shear rates of 0.05 and 5 s1 in heating and cooling cycles. As shown in Figure 8.4, the transition temperature, as defined earlier, is found 72 and 60
C at the low and high shear rate, respectively. The lower transition temperature indicated that the exfoliation was more preferable at the high shear rate. At 136
C, viscosity at the high shear rate is seen to be significantly lower than that at the low shear rate (Figure 8.4). The lower viscosity at the high shear rate was thought to be due to a stronger shear-induced orientation, leading to the lower resistance to flow.

FIGURE 8.4 Temperature dependence of the steady shear viscosity at shear rates of 0.05 and 5 s1 for HTPB4200–clay (5 wt.%) composite [8].

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The results obtained from steady and oscillatory flows were also compared. It was interesting that the oscillatory flow did not show any evidence of the shear-induced orientation leading to a reduction of viscosity. According to authors, at small strain amplitudes there is insufficient driving force to induce the clay orientation. Therefore, the steady flow is preferable in order to induce the shear-induced clay orientation leading to the reduction of viscosity. Zhu et al. [20] studied rheological behavior of liquid BR–clay composites by varying functional groups of liquid rubber and organic modifiers in clay. The materials used in their study were liquid PB oligomer, HTPB, and carboxylterminated polybutadiene (CTPB), organoclay modified with octadecyltrimethylammonium chloride (C18A) and dioctadecyldimethylammonium chloride (D18A). The composites were prepared by solution mixing of liquid rubber and clay at a loading of 10 wt.% at room temperature. Rheological test was carried out according to similar procedures as in Ref. [8]. Viscosity of HTPB–C18A, HTPB–D18A, and CTPB–C18A composites upon heating from a temperature of 26
C increased with temperature up to their transition temperature, then slightly decreased at higher temperature up to 116
C. The increase of viscosity was due to the rubber–clay interaction and exfoliation caused by the chemical reaction between functional groups in liquid rubber and clay. After cooling from 116 to 26
C, the viscosity of these composites was significantly higher than the viscosity at the start of heating, indicating the irreversible exfoliation generated during heating. However, for CTPB–D18A composite, the increase in viscosity during heating, and the difference in viscosity before heating and after cooling were not observed. This behavior was an indication of the absence of exfoliation during heating. However, the results obtained by XRD and TEM for this composite revealed that the exfoliation and good dispersion of clay were already achieved before heating in the rheometer. Hence, the organoclay modified with D18A (dioctadecyldimethylammonium chloride) containing two long alkyl chains was fully exfoliated in the carboxyl terminated polybutadiene liquid rubber by the solution mixing. 8.2.4 Acrylonitrile Butadiene Rubber (NBR)–Clay Nanocomposites Kim et al. [21] studied structures and properties of NBR–clay composites. Organic modification of Na-montmorillonite, supplied by Kunimine Co., by octyl amine CH3(CH2)7NH2 (C8), dodecyl amine CH3(CH2)11NH2 (C12) and octadecyl amine CH3(CH2)17NH2 (C18) were used. NBR with 29% of acrylonitrile (AN) content was mixed with the organoclays in an internal mixer (Brabender Plasticorder). Mooney viscosity ML1þ4 of the composites at 100
C was measured. A value of ML1þ4 as a function of clay content was dependent on the type of the organic modifiers, as shown in Figure 8.5. ML1þ4 of the composite using octyl amine (C8) modified clay increased with the clay content to 10 phr, which is a general observation with filler loading. However, ML1þ4 of the composite using dodecyl amine (C12) modified clay did not change with the clay content. Furthermore, the composite with the clay modified with octadecyl amine (C18) displayed a decrease in ML1þ4 with the clay content to 10 phr. The authors explained their observations by suggesting that an

RHEOLOGICAL BEHAVIOR OF RUBBER–CLAY NANOCOMPOSITES

C8-MMT C12-MMT C18-MMT

50 M1+4

251

45

40 0

2

6 4 Clay content (phr)

8

10

FIGURE 8.5 ML1þ4 at 100
C as a function of clay loading containing various organic modifiers [21].

organic modifier for clay may act as a plasticizer. However, the ability of organoclay to act as a plasticizer is dependent on the chain length of the organic modifier. The organic modifier with a longer chain length such as octadecyl amine (C18) acts as a plasticizer causing a reduction of ML1þ4 of the composite in spite of an increase in clay content. On the other hand, the organic modifier with a shorter chain length such as octyl amine (C8) does not act as a plasticizer at all. Therefore, a chain length of an organic modifier in clay is an important factor in determining the rheological behavior of NBR–clay composite. Accordingly, depending on the type of organic modifier in the clay, viscosity of NBR–clay composite can be lowered in comparison with NBR matrix. Sadhu and Bhowmick [7] studied rheological behavior of NBR–clay composites. The composite preparation and experimental methods were similar to those used in SBR–clay composite discussed earlier. NBR rubbers with different AN contents of 19, 35, and 50% were used. The decrease in viscosity by adding clay was found, as shown in Figure 8.6. In particular, viscosity of NBR (34% of acrylonitrile) composites containing 2, 4, and 8 phr of the organoclay was lower than that of unfilled NBR. However, viscosity of the composite with 8 phr of organoclay was higher than that of composite with 4 phr of organoclay. The authors explained such a behavior as a result of the orientation of clay platelets in a capillary flow at high shear rates and the slippage of NBR chains on the clay surface. The increase in viscosity at the organoclay concentration of 8 phr was explained by a formation of clay agglomerates occurring at this high concentration of clay. Another observation seen from Figure 8.6 is the difference in viscosity of the composites filled with the pristine Na–montmorillonite and the organically modified montmorillonite. Viscosity of the NBR–pristine clay composite at a loading of 4 phr was significantly higher than that of the NBR–organoclay composite. A more smooth extrudate surface and a lower die swell for the NBR–organoclay composite were also observed compared to those of NBR–pristine clay composite at a loading of 4 phr.

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FIGURE 8.6 The shear viscosity versus the shear rate of NBR containing 34% of acrylonitrile group and its clay composites at 130
C. NBRN4 corresponds to the composite with 4 phr of pristine clay, while NBROC2, NBROC4, and NBROC8 correspond to the composite with 2, 4, and 8 phr of the organoclay, respectively [7].

A significant reduction of viscosity by adding the organoclay was observed in NBR with 34 and 50% of AN content. However, the composite containing NBR with 19% of AN content did not show a decrease in viscosity by adding the organoclay. Clearly, the level of polarity of matrix may also play an important role in determining the orientation of the organoclay in the composite during its flow at high shear rates. Ibarra et al. [22] studied structures and properties of NBR–clay composites. Carboxylated nitrile rubber (XNBR) containing 7 wt.% of carboxylic group was mixed with organoclays at a loading of 7 phr in a “cylinder mixer.” The clays used were montmorillonite modified with hydrogenated ammonium dimethyl-ditallow chloride (Cloisite 15A, Southern Clay) and sodium bentonite supplied by Tolsa SA treated with octadecylammonium (BCA-ODA). Rheological measurements were performed using a capillary rheometer (Rheograph 2003) and a Mooney viscometer (MV-2000E) at 90
C. Viscosity ML2 þ 4 of BCA-ODA, Cloisite 15A composites and unfilled XNBR were 3.3, 3.6, and 2.7, respectively. Mooney viscosity of XNBR–Cloisite 15A composite is seen to have the highest value. However, viscosity measured by the capillary rheometer showed the higher viscosity for XNBR–BCAODA composite compared to that of XNBR–Cloisite 15A in the whole range of shear rates. The power-law indexes (n) of BCA-ODA, Cloisite 15A composites, and unfilled XNBR were 0.375, 0.357, and 0.339, respectively. The consistency indexes were 210, 157, and 105 kPa sn for BCA-ODA, Cloisite 15A composites, and unfilled XNBR, respectively. The lowest die swell was observed for the composite containing BCA-ODA. The XRD data showed the exfoliation and intercalation for the XNBR–BCA-ODA and XNBR–Cloisite 15A composite, respectively. Clearly, the

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253

viscosity behavior of XNBR–organoclay reported in Ref. [22] were different from that of NBR–organoclay in Ref. [7]. 8.2.5 Ethylene Propylene Rubber–Clay Nanocomposites Austin and Kontopoulou [23] studied structures and properties of ethylene–propylene rubber (EPR)–clay composites. Maleated EPR containing 0.5–1.0 wt.% of maleic anhydride and organoclay modified with dimethyl dialkylammonium halide containing 70% of C18 (I.44PA, Nanocor) were used. Rheological properties in frequency sweep mode were obtained using an oscillatory rheometer with parallel plates (Rheologica ViscoTech) at a temperature of 170
C. A significant increase in the complex viscosity and storage modulus was observed by adding clay up to 20 wt.%. In particular, a Newtonian behavior at low frequencies was observed at a clay loading of 5 wt.%, while such a behavior was not present at a loading of 10 wt.%. The concentration of 5% of clay was insufficient for creating a clay percolation as confirmed by a morphological analysis using TEM. Hence, the increase in viscosity by adding the organoclay up to 5 wt.% was due to the rubber–clay interaction alone and not by the clay–clay interaction. The amount of bound rubber was measured, showing the evidence of the rubber–clay interaction. At high loading of clay, both the clay–clay interaction as well as the rubber–clay interaction was evident significantly affecting the complex viscosity and storage modulus. The rheological percolation of clay occurred between the clay loadings of 5 and 10 wt.%. Mohammadpour and Katbab [24] studied structures and properties of EPDM rubber–clay nanocomposites containing maleated EPDM as a compatibilizer between EPDM and clay. The rubbers used were EPDMs with ML1þ4 at 125
C of 25 and 85, and the maleated EPDM containing 0.4% maleic anhydride. The clay modified with dimethyl dehydrogenated tallow ammonium (Dellite 67G, Laviosa Chemica Nineraria) was used. The master batch of organoclay at a loading of 5% and maleated EPDM at loadings of 5 and 15% were prepared using a Banbury mixer. This master batch was compounded with EPDM in a two-roll mill. Rheological behavior in frequency sweep mode was obtained using the oscillatory rheometer with parallel plates (Paar Physica US 200) at a temperature of 180
C and a strain amplitude of 1%. In the low frequency region, the storage modulus of the EPDM–clay composites containing 5% of clay and 5% of compatibilizer was higher than that of unfilled EPDM. The increase in the storage modulus was due to the intercalation of clay as verified by XRD results. In particular, the EPDM–clay composite containing 15% of maleated EPDM exhibited a higher storage modulus at low frequencies than that of the composite containing 5% of maleated EPDM. The increase was due to the enhanced polarity of EPDM and its better compatibility with clay. The increase in the storage modulus with concentration of compatibilizer was higher for EPDM of high viscosity. This observation was in line with the XRD results. A complete exfoliation was found in the high viscosity EPDM–clay composite, while a partial exfoliation in the low viscosity EPDM–clay composite. As a frequency increased, the storage modulus of the composites approached to that of the unfilled EPDM. This was

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ascribed to the reduced chain entanglement caused by the dispersed clay, leading to a lower increase of the storage modulus of composite at high frequencies. Tan and Isayev [25,26] studied properties of EPDM rubber–clay nanocomposites processed in ultrasonic extruder in comparison with carbon black and silica composites. EPDM was mixed with montmorillonite modified with dimethyl dehydrogenated tallow amine (Cloisite 15A, Southern Clay) at loadings of 5 and 10 phr in an internal mixer, and processed in a single screw extruder [26]. The composites were ultrasonically treated during processing in the extruder [25]. Rheological properties in frequency sweep were obtained using an Advanced Polymer Analyzer (APA 2000) at a temperature of 100
C and a strain amplitude of 7.1%. The storage modulus increased with the loading of clay over all the frequency range. The enhanced reinforcement by adding clay was also verified by measurements of the mechanical properties, that is, the increased tensile strength, elongation at break, and modulus at 100% strain. The storage modulus at the clay loading of 10 phr was higher than that of carbon black (V1391, Cabot) filled composite, while it was lower than that of silica (Hi-Sil132, PPG) filled composite at the same loading. After the composites were treated by ultrasound at the amplitude of 6.5 mm, the storage and loss moduli as well as the complex dynamic viscosity of the composites were slightly lower than those of untreated composites, due to rubber chains scission by ultrasound. However, mechanical properties of composites were increased or unaffected by ultrasonic treatment. Susteric and Kos [9] studied rheological behavior of EPDM–clay composites using experimental data and model based on statistical mechanics of random networks. EPDM was mixed with organically modified montmorillonite with ditallow dimethyl amine (Dellite 67G, Laviosa Chimica Mineraria) at loadings of 5, 10, 20, and 30 phr in an internal mixer. Rheological behavior in strain sweep mode was obtained using RPA 2000 at a frequency of 0.3 Hz in the temperature range from 40 to 100
C. Based on rheological data and model it was shown that the storage modulus increased and the loss modulus decreased with organoclay loadings. 8.2.6 Fluoroelastomer–Clay Nanocomposites Kader et al. [10] studied rheological behavior of fluoroelastomer–clay composites in comparison with carbon black and silica composites. Fluoroelastomer (FKM, Viton A-500), Na-montmorillonite and organophilic montmorillonite modified with quaternary ammonium salt of dimethyl hydrogenated tallow (Cloisite 15A, Southern Clay) were used. Fluoroelastomer was mixed with clays at loadings up to 20 wt.% in a two-roll mill. Rheological test was performed using a capillary rheometer (Monsanto Processability Tester) at temperatures of 130, 150, and 170
C. The viscosity as a function of the shear rate for the composites filled with the organoclay and the pristine clay at 150
C are presented in Figure 8.7a and b, respectively. The power-law index (n) and consistency index (k) calculated by the fitting the power-law equation to the viscosity plots at the temperature of 150
C are presented in Table 8.1. As shown in Figure 8.7a, viscosity of all the composites containing the organoclay is lower than that of unfilled fluoroelastomer. As seen from Table 8.1, the organoclay-filled composites show the higher power-law index and lower consistency index than

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255

FIGURE 8.7 Shear viscosity versus shear rate for fluoroelastomer and its clay composites at 150
C and various contents of the organoclay (a) and the pristine clay (b) [10].

TABLE 8.1 Power-Law Index (n) and Consistency Index (k) for Fluoroelastomer–Clay Composites at 150
C No Filler

Composites Filled with Organoclay

Composites Filled with Pristine Clay

Clay wt.% 0 2.5 5 7.5 10 15 20 2.5 5 15 20 n 0.12 0.20 0.19 0.20 0.20 0.20 0.18 0.12 0.14 0.13 0.12 k (Pa sn) 26.4 13.6 14.2 12.2 12.1 11.5 13.7 25.2 22.0 24.8 27.8

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unfilled fluoroelastomer. Such a behavior was ascribed to the exfoliation and subsequent flow-induced orientation of the clay platelets as well as the slippage of rubber chains on the clay surface. The plasticizing behavior of the organic modifier for clay was also mentioned as one of reasons of the viscosity reduction. On the other hand, as shown in Figure 8.7b, the pristine montmorillonite-filled fluoroelastomers does not show any remarkable difference in viscosity compared to unfilled fluoroelastomer. The power-law index and consistency index of unfilled fluoroelastomer and fluoroelastomer–pristine clay composites in Table 8.1 are close to each other. In contrast, viscosity of the carbon black (N550, Kumho Petrochemical) and silica (Ultrasil VN3, Degussa) filled fluoroelastomer composites at a loading of 10 wt.%, was higher than that of the unfilled fluoroelastomer. The decrease in viscosity with loading of organoclay was the highest at low shear rates, while it was less at high shear rates. At low shear rates, the shear flow resistance was dominated by the state of dispersion of clay platelets and their orientation along the flow direction, while at the high shear rate it was dominated by the viscoelastic characteristics of rubber matrix. The activation energy of viscous flow at particular shear rates (Eg_ ) was calculated using the Arrhenius–Frenkel–Eyring equation:   Eg_ ha ¼ B exp ð8:5Þ RT where ha is the viscosity at a particular shear rate, B is the pre-exponential factor, R is the gas constant, and T is the absolute temperature. The activation energy for the fluoroelastomer–organoclay composites was higher ranging from 5 to 22 kJ/mol compared to that of unfilled fluoroelastomer and fluoroelastomer–pristine clay composites. This indicated higher sensitivity of viscosity to temperature variation for the fluoroelastomer–organoclay composites. The value of composites die swell at a temperature of 150
C decreased from 1.3 to 0.95 with loadings of organoclay up to 20 wt.%. The extrudate surface was wavy for unfilled fluoroelastomer and composite at 5 wt.% of organoclay at high shear rates, but upon adding organoclay at a loading of 10 wt.%, it became smooth. The surface smoothness was better for the organoclayfilled composites than for the pristine clay filled composites. Lakshminarayanan et al. [27] studied structures and properties of fluoroelastomer–clay composites. The clays used were pristine Na-montmorillonite (Cloisite Na, Southern Clay), montmorillonite modified with di(hydrogenated tallow-alkyl) dimethyl ammonium (Cloisite 15A, Southern Clay), di(hydrogenated tallow-alkyl) dimethyl ammonium (Cloisite 20A, Southern Clay), with tallow alkyl methyl di(2hydroxyethyl)ammonium (Cloisite 30B, Southern Clay) and di(hydrogenated tallowalkyl)methyl ammonium (Cloisite 93A, Southern clay). These clays were compounded with fluoroelastomer (Dyneon FPO 3741) containing 69.5% of fluorine using an internal mixer (Haake Rheomix Series 600) at a temperature of 75
C. The clay concentrations were 2.5, 5.0, and 10 phr. Rheological properties in frequency sweep mode were obtained using an oscillatory rheometer with parallel plates (Rheometrics RMS 800) at a temperature of 180
C. For the composites containing Cloisite 20A and 15A, the storage modulus increased with loading of

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257

organoclays, but the slope of the storage modulus versus frequency at low frequencies decreased with loading of organoclay. The latter indicated a transition from the liquid-like to pseudo solid-like behavior. At the same time, for the composites containing Cloisite 93A, 30B, and pristine clay, the storage modulus and slope of the storage modulus versus frequency did not change with the loading of clays. This was due to the absence of intercalation as indicated by XRD results. In particular, the interlayer spacing of the composite containing Cloisite 30B decreased with clay loading, possibly due to the loss of organic modifier. In contrast, the interlayer spacing of the composite containing Cloisite 20A increased compared to that of the composite containing pristine clay, leading to the increase of the storage modulus by adding the clay. Interestingly, the composites containing Cloisite 15A did not show intercalation at all as indicated by XRD study, although the storage modulus of these composites increased with the loading of clay. The authors ascribed such a behavior to the competing effects in motion of the organic modifier and rubber chains. In particular, the unbound organic modifier in the clay diffused out during compounding, leading to the decrease of interlayer spacing. At the same time, as shown in TEM images, rubber chains are intercalated to the interlayer of organoclay, leading to the increase of interlayer spacing. Thus, due to these competing effects, a significant shift of peak in XRD data with the loading of clay was not observed. The observed increase of the storage modulus with the loading of clay was apparently due to the penetration of rubber chains to the interlayer of organoclay. 8.2.7 Poly(isobutylene-co-para-methylstyrene) (BIMS) Rubber–Clay Nanocomposites Maiti et al. [28] studied structures and properties of BIMS–clay composites. The materials used were BIMS rubber containing 7.7 wt.% of paramethyl and 1.2 wt.% of bromine (BIMS7745, Exxon Mobil) and organoclay prepared by treating Namontmorillonite, supplied by Southern Clay, with octadecyl amine (C18H37NH2). The BIMS rubber–clay composites at a loading of 4 wt.% were prepared by a solution mixing. Viscosity was measured using a capillary rheometer (Monsanto Processability Tester) at a temperature of 130
C. Viscosity of unfilled BIMS and BIMS–organoclay composite were found to be similar, while viscosity of the BIMS–pristine clay composite was slightly lower. These results were not correlated with the result from XRD study, showing intercalation for BIMS–pristine clay composite and exfoliation for BIMS–organoclay composite. The power-law index of 0.09, 0.13, and 0.12, and the consistency index of 86, 64, and 65 Pa sn were obtained for BIMS–organoclay, BIMS–pristine clay composites, and unfilled BIMS, respectively. The die swell of BIMS–organoclay composite was lower than that of BIMS–pristine clay composite. 8.2.8 Poly(ethylene-co-vinylacetate) (EVA) Rubber–Clay Nanocomposites Riva et al. [29] studied structures and properties of EVA rubber–clay nanocomposites. The EVA rubber contained 19 wt.% of vinylacetate. This rubber is used for

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shoes and telecommunication cable. The organophilic clays were prepared by modification of synthetic fluorohectorite (Somasif ME100, Co-Op Chemical) with octadecylammonium (ME100/ODA) and aminododecanoic acid (ME100/ADA), montmorillonite modified with methyl tallow bis-2-hydroxyethyl ammonium (Cloisite 30B, Southern Clay) and dimethyl dehydrogenated tallow ammonium (Cloisite 6A, Southern Clay). The EVA rubber and 10 wt.% of clay were compounded in the internal mixer (Brabender). Rheological properties in frequency sweep mode were obtained using an oscillating rheometer with parallel plates (Dynamic Analyzer Rheometer RDA II) at a temperature of 120
C and a strain amplitude of 0.05%. The storage modulus of the composite containing Cloisite 30B showed the highest value in all the frequency range. The storage modulus of the composites containing ME100/ADA and Cloisite 6A was similar to that of unfilled EVA rubber. The storage modulus of EVA composite containing ME100/ODA had an intermediate storage modulus. The difference in the storage modulus was due to the extent of clay exfoliation and intercalation. The exfoliation of Cloisite 30B and ME100/ODA was confirmed by the XRD and TEM studies. The intercalation without exfoliation was observed in the composite containing Cloisite 6A. On the other hand, no intercalation or exfoliation of clay was seen in the composite containing ME100/ADA. Therefore, the exfoliation and intercalation play an important role in enhancing the storage modulus of the EVA composites and dependent on the compatibility between organoclay and rubber matrix. Prasad et al. [11] studied rheological behavior of EVA–clay composites under shear and extensional flow. EVA containing 9 wt.% of vinyl acetate and Namontmorillonite modified with dimethyl dehydrogenated tallow amine (Cloisite 15A, Southern Clay) were used. EVA and organoclay were mixed in a twinscrew extruder at clay loadings of 2.5, 5.0, and 7.5 wt.%. Rheological properties were obtained using a rheometer with parallel plates in both oscillatory and steady shear modes (ARES) at a temperature of 130
C. In an oscillatory mode, a frequency of 10 rad/s and strain amplitudes of 5, 4, 3, and 2% at clay loadings of 0, 2.5, 5, and 7.5 wt.% were applied, respectively. The plots of the storage modulus versus a frequency are presented in Figure 8.8a. The storage modulus increases with the loading of clay. The slope of G0 versus v in the low frequency region was found to decrease with the loading of clay, indicating that percolation was not achieved even at a loading of 7.5 wt.%, though rheological behavior showed a tendency toward the pseudo solid-like behavior. Steady shear viscosity at low shear rates increased with the loading of clay, due to the increased interaction between polymer and clay. The first normal stress difference (N1) as a function of the shear stress showed a linear correlation without the dependency of clay loading. This independency was due to the preferential orientation of clay in the flow direction leading to the reduction of the contribution of clay to the elasticity of composite. Uniaxial extensional rheological behavior was obtained using rheometrics melt extensional (RME) rheometer at a temperature of 130
C and strain rates of 0.1, 0.5, and 1 s1. The plots of extensional viscosity versus time are presented in Figure 8.8b. Unfilled EVA and all the composites show an increase of the extensional viscosity with time and strain rate. The extensional viscosity of the unfilled EVA and composites at loadings of

RHEOLOGICAL BEHAVIOR OF RUBBER–CLAY NANOCOMPOSITES

259

FIGURE 8.8 The storage modulus versus frequency (a), extensional viscosity versus time at various strain rates (b) of EVA–organoclay composites [11].

2.5 and 5 wt.% is similar. However, the extensional viscosity of the composite containing 7.5 wt.% of clay is slightly higher than that of the other composites, indicating the stronger interaction between clay particles. 8.2.9 Polyepichlorohydrin Rubber–Clay Nanocomposites Lim et al. [12] studied rheological behavior of polyepichlorohydrin rubber–clay nanocomposite. PECH rubber exhibiting a high resistance to ozone, oil, heat, and weathering, and organoclay modified with dimethyl hydrogenated tallow quaternary ammonium methylsulfate (Cloisite 25A, Southern Clay) were used. PECH rubber

260

RHEOLOGY OF RUBBER–CLAY NANOCOMPOSITES

FIGURE 8.9 Steady shear viscosity versus shear rate for PECH rubber and PECH rubber– clay composites at various clay loadings at 200
C. The line represents the curve fitting using the Carreau model [12].

was solution-mixed with organoclay at loadings of 2, 5, and 10 wt.%. A steady shear viscosity was measured using a rotational rheometer with parallel plates (Physica MC 120) at a temperature of 200
C. As shown in Figure 8.9, a significant increase in viscosity and shear thinning behavior are seen with the loading of clay. Viscosity as a function of shear rate was fitted to the Carreau model: h¼

h0

ð8:6Þ

½1 þ ð_glÞ2 ð1  nÞ=2

where, h0 is the zero-shear viscosity, g_ is the shear rate, l is the relaxation time. The zero-shear viscosity and relaxation time are presented in Table 8.2. An increase of the zero-shear viscosity with the clay concentration is clearly seen. This is due to the enhancement of the rubber–clay interaction and the dispersion of clay. Interestingly, the composite with 10 wt.% of clay showed the strong viscosity drop at high shear rates, with its viscosity being even lower than that of the composite containing 5 wt.% TABLE 8.2 Zero-Shear Viscosity and Relaxation Time in Carreau Model for the PECH and PECH Rubber–Clay Nanocomposites

PECH PECH–clay 2 wt.% PECH–clay 5 wt.% PECH–clay 10 wt.%

h0 (Pa s)  10  5

l (s)

0.44 0.57 0.87 1.65

3.27 2.49 3.58 19.36

RHEOLOGICAL BEHAVIOR OF RUBBER–CLAY NANOCOMPOSITES

261

of clay. Evidently, this behavior was due to the alignment and orientation of clay particles along the flow direction. 8.2.10 Thermoplastic Polyurethane (TPU)–Clay Nanocomposites Berta et al. [30] studied structures and rheological properties of TPU–clay composites. TPU was synthesized using polyol(ethylene oxide-tipped polypropyleneoxide polyether), catalyst (triethylenediamine), and isocyanate (methylenediphenylendiisocyanate (MDI)). Two composites were prepared by mixing polyol and organoclay (Cloisite 30B, Southern Clay) for a period of 5 min (short mixing) and 1 h (long mixing). Then, these blends were mixed with catalyst and isocyanate. A hard block content in TPU was 30%. An organoclay content in composite was 2.5 wt.%. Rheological properties in frequency sweep mode were obtained using DMTA (Polymer MK II) at room temperature and a strain amplitude of 0.6% under N2 atmosphere. Due to intercalation and exfoliation of clay, the storage modulus of the composite mixed for the long time was significantly higher than that of pure PU. This was supported by XRD and TEM studies. Therefore, it was concluded that the higher storage modulus was caused by the stronger interaction between rubber and clay due to the better dispersion. In contrast, the composite mixed for the short time showed no intercalation and poor dispersion as indicated by XRD and TEM data, respectively. The storage modulus of this composite was lower than that of the composite mixed for the long time and even lower than that of the pure PU. Mishra et al. [31] studied structures and rheological properties of TPU–clay composites. The polyester-based TPU with a trade name of Desmopan KU 8600E prepared from MDI, polytetramethylene glycol, and 1,4-butanediol were used. Two clays such as synthetic layered silicate (Laponite RD, Rockwood) and montmorillonite modified with dimethyl dehydrogenated tallow amine (Cloisite 20A, Southern Clay) were used. Laponite RD was organically modified with dodecylamine hydrochloride (O-Laponite RD). TPU–clay composites were prepared by adding the organoclays at loadings of 1, 3, 5, 10 wt.% in TPU-THF solution followed by THF evaporation. Rheological test was carried out using a dynamic mechanical analyzer in a tension mode (DMA 2980, TA) at a frequency of 1 Hz, a strain amplitude of 0.1% and a temperature range of 60–100
C. The storage modulus of TPU–Cloisite 20A composites continuously increased with the loading of clay. However, the storage modulus of TPU–O-Laponite RD composites showed a maximum at a loading of 3 wt.%. The storage modulus at the loading of 10 wt.% was even lower than that of pure TPU. This behavior was ascribed to the reagglomeration of the O-Laponite RD at concentrations more than 3 wt.%, due to its higher hydrophilicity and lower content of the organic modifier than those of Cloisite 20A. Also, rheological properties were measured using an oscillatory rheometer (RPA 2000) at a temperature of 140 and 170
C and a strain amplitude of 1%. At a temperature of 140
C when hard domains in TPU matrix are ordered, tan d at 3.33 Hz was 0.31–0.49 for all the composites. However, at a temperature of 170
C when hard domains in TPU matrix are melted, tan d at 3.33 Hz was 1.36–3.87 for all the composites. At 140
C and 3.33 Hz, the increase in the storage modulus by adding clay was more pronounced in case of TPU–O-Laponite

262

RHEOLOGY OF RUBBER–CLAY NANOCOMPOSITES

RD composites. The latter was due to the preferential association of O-Laponite RD with the hard segment in TPU matrix. It should be noted that Cloisite 20A preferentially associates with the soft segment in TPU matrix [13]. However, at 170
C and 3.33 Hz, the storage modulus of TPU–O-Laponite RD composites decreased with the loading of clay, due to its ability for a strong agglomeration. This behavior was not observed for the TPU–Cloisite 20A composites, because Cloisite 20A has a higher aspect ratio and a lower cohesive force leading to the less agglomeration than O-Laponite RD. 8.2.11 Styrene–Ethylene–Butylene–Styrene (SEBS) Block Copolymer–Clay Nanocomposites Chen et al. [32] studied structures and rheological properties of SEBS–clay composites containing compatibilizer. SEBS with 29 wt.% of styrene, SEBS grafted with maleic anhydride (SEBS-g-MA), polypropylene grafted with maleic anhydride (PPg-MA), and montmorillonite modified with dimethyl dehydrogenated tallow amine (Cloisite 20A, Southern Clay) were used. SEBS and organoclay at loadings of 2, 5, and 7 wt.% were mixed with 20 phr of compatibilizer (SEBS-g-MA or PP-g-MA) in an internal mixer. Rheological properties were measured in the steady shear flow using a rheometer (AR 2000, TA) at a temperature of 180
C. Viscosity increased with the loading of clay for all the composites. The composite containing PP-g-MA and SEBS-g-MA showed significantly higher viscosity than that without a compatibilizer. As indicated by XRD studies, the increase of the interlayer spacing of the composite containing SEBS-g-MA was more than that without a compatibilizer. Relatively thin layered silicates were observed in TEM images for the composites containing compatibilizers. Thus, the observed increase in viscosity of composites containing compatibilizers was due to the enhanced interaction between polymer and clay. Carastan et al. [14] studied rheological behavior of SEBS–clay and SBS–clay composites. One SBS and several SEBS were used in SEBS-30 containing 30 wt.% of styrene, SEBS-13/29 containing 13 wt.% of styrene and 29 wt.% of diblock, SEBS30/70 containing 30 wt.% of styrene and 70 wt.% of diblock, and SEBS-MA containing 30 wt.% of styrene and 1–2 wt.% of maleic anhydride. Clay was montmorillonite modified with dimethyl dehydrogenated tallow amine (Cloisite 15A, Southern Clay). The composites at the clay loading of 5 wt.% were prepared by a solution mixing, melt mixing in an internal mixer and a combination of these two mixing. Rheological behavior in frequency sweep mode was obtained using an oscillatory rheometer with parallel plates (ARES TA) at a temperature ranging from 130 to 220
C and a strain amplitude of 1%. At the temperature corresponding to disordered state of the block copolymer (200
C), the storage modulus and dynamic viscosity of the composites were higher than those of the unfilled block copolymer. In particular, the SBS and SEBS-13/29 composites prepared by solution mixing showed the highest storage modulus and dynamic viscosity with a strong pseudo solid-like behavior. The composites prepared by melt mixing showed the lowest values. The difference in viscosity due to the preparation method was ascribed to the state of

GENERAL REMARKS ON RHEOLOGY OF RUBBER–CLAY NANOCOMPOSITES

263

dispersion observed by the optical microscopy and TEM images. In particular, the level of exfoliation of clay was quantified from the analysis of TEM images [33], showing 58% for solution mixing, 48% for combination of solution and melt mixing, and 10% for melt mixing. Also, the presence of an interconnected network structure of clay in the solution mixed composite was considered as one of the reasons of the higher storage modulus and viscosity in comparison with the composite prepared by melt mixing. However, the storage modulus and viscosity at 200
C of SEBS-30/ clay, SEBS-30/70/clay, SEBS-MA/clay composites were not dependent on the mixing method. This was due to the presence of cylindrical ordered state in matrices, instead of the disordered state, hindering the effect of mixing method on the rheological behavior. It was found that the addition of clay has a little effect on the rheological properties for a lamellar or cylindrically ordered block copolymer, because these structures already exhibit the pseudo solid-like behavior. Among the solution prepared composites, the SEBS-MA composite showed the highest storage modulus without presence of the terminal region, due to the better dispersion and strong interaction between clay and polymer created by the presence of maleic anhydride.

8.3 GENERAL REMARKS ON RHEOLOGY OF RUBBER–CLAY NANOCOMPOSITES The summary of the investigations on rheological behavior of rubber–clay nanocomposites is presented in Table 8.3. Rheological behavior of rubber–clay nanocomposites was obtained using steady shear flow [7,8,10–12,16,20–22,28,32], oscillatory shear flow [6,8,11,13,14,17–20,23–27,29–31,33] and extensional flow [11]. By fitting the power-law equation to steady shear viscosity data, the power-law and consistency index were obtained [6,7,10,16,17,22,28]. Typically, the power-law index decreased with the loading of clay. The percolation threshold of clay was obtained by observing variations of slopes of the storage modulus in the low frequency region with clay loading [6,11,17,23]. Based on these data, the percolation thresholds of organoclay were 4, above 5, and above 7.5 wt.% in IR, EPR, and EVA matrices, respectively. The activation energy of viscous flow was obtained using Arrhenius–Frenkel–Eyring equation. A higher value of the activation energy was obtained for fluoroelastomer–organoclay composites compared to unfilled and pristine clay filled fluoroelastomer [10]. By fitting Carreau model to the viscosity data, zero-shear viscosity and relaxation time were obtained. Polyepichlorohydrin rubber–organoclay composites show the significant increase of these parameters with the loading of clay [12]. The calculation of the aspect ratio of clay in IR composites was attempted using two models, showing the values of approximately 282 [6] and 62 [17] with the experimental value being 110. The extent of Payne effect in strain sweep experiments was used to evaluate the presence of clay networks in the composites with matrices of IR [18], ENR [18], SBR [19], and EPR [23]. Also, the linear region of composite dynamic behavior was verified by plotting the storage modulus versus a strain amplitude observed in the

264

(Rubbers: BR, NBR19, NBR50)

Rubbers/organoclay 4 phr Extrudate image

Unusual behavior was seen for NBR and SBR composites

Die swell versus g_

(Rubbers: SBR, NBR34)

Power-law index (n) Consistency index (k)

h versus g_

Rubbers/Cloisite Na 4 phr

Rubbers/organoclay 2, 4, 6, 8 phr

G0 versus strain

SBR/TESPT/fibrillar silicate 40 phr

(Rubbers: IR, ENR25, ENR50)

Rubbers/silica 10 phr

Rubbers/organoclay (C3) 10, 20, 30 phr

The extent of Payne effect by the addition of TESPT and additional processing

Comparison between silica and organoclay

tan d versus strain

Rubbers/organoclay (C1) 10, 20, 30 phr

Rubbers/organoclay (C2) 10, 20, 30 phr

The extent of Payne effect in comparison to different composites

G0 versus strain

Rubbers/Na-montmorillonite 10, 20, 30 phr

Power-law index (n) Slope of G0 in the terminal flow region to obtain the percolation threshold Relative viscosity to obtain the percolation threshold Evaluation of the aspect ratio using model equations Alignment effect of clay on G0

G0 versus strain G0 , G00 , h0 versus v

Power-law index (n)

h versus g_

Rubbers/bentonite 1, 1.5, 2, 2.5 phr Rubbers/fluorohectorite 1, 1.5, 2, 2.5 phr (Rubbers: NR latex, XSBR latex)

IR/Cloisite 10A 1, 2, 3, 4, 5, 7, 9 wt.% IR/Cloisite Na 5 wt.%

Remarks

Data Obtained

Materials

TABLE 8.3 The Summary of Studies on Rheological Behavior of Rubber–Clay Nanocomposites

[7]

[19]

[18]

[6,17]

[16]

References

265

EPDM/silica 10, 20, 30 phr

EPDM/carbon black 10, 20, 30 phr

EPDM/Cloisite 15A 5, 10 phr

EPDM/compatibilizer/Dellite 67G 5%

EPR/I.44PA 2, 5, 10 wt.%

G0 , G00 , |h |, tan d versus v



G0 , G00 versus v

G0 , G00 , |h |, tan d versus v



G0 versus strain

Effect of ultrasound on the rheological properties

Comparison between clay, silica and carbon black

Effect of compatibilizer and viscosity of EPDM on the storage modulus

The extent of Payne effect by adding clay

Slope of G0 in terminal flow region to obtain percolation threshold

Consistency index (k)

h versus g_

XNBR/organoclay 7 phr Extrudate image

Power-law index (n)

Viscosity dependency on different organoclays

Viscosity dependency on heating and cooling in comparison to different composites

Comparison between steady shear and oscillatory shear flow

Viscosity dependency on heating and cooling at different shear rate

ML2 þ 4

ML1 þ 4 versus clay loading

h versus temperature

G versus time

0

G versus temperature

0

h versus temperature

XNBR/Cloisite 15A 7 phr

(Organoclays: C8, C12, C18)

NBR/organoclays 1, 2, 5, 10 phr

(Liquid BRs: PB, HTPB, CTPB)

Liquid BRs/organoclay (D18A) 10 wt.%

Liquid BRs/organoclay (C18A) 10 wt.%

Liquid PB/organoclay 5 wt.%

Liquid HTPB/organoclay 5 wt.%

(continued)

[25,26]

[24]

[23]

[22]

[21]

[20]

[8]

266

EVA/Cloisite 15A 2.5, 5, 7.5 wt.%

(organoclays: Cloisite 6A, Cloisite 30B, Somasif ME100/ODA, Somasif ME100/ADA)

EVA/organoclays 10%

Extensional viscosity (hE) Melt strength of extruded filament at various draw down ratios (DDRs)

N1 versus s

Dependency of elasticity on different organoclays

h versus g_ hE versus time

Slope of G0 in terminal flow region to obtain percolation threshold

G0 , G00 versus v

G versus v

Storage modulus dependency on different organoclays

Consistency index (k)

Die swell versus g_

BIMS/Na-montmorillonite 4 phr 0

Power-law index (n)

Storage modulus dependency on different organoclays

h versus g_

G0 , tan d versus v

Activation energy (Eg_ ) Comparison between clay, silica and carbon black

BIMS/organoclay 4 phr

(clays: Cloisite Na, 15A, 20A, 30B, 93A)

Fluoroelastomer/clays 2.5, 5, 10 phr

Extrudate image

Consistency index (k) Die swell

Fluoroelastomer/Na-motmorillonite 2.5, 5, 10, 15 phr

00

Power law index (n)

00

tw versus g_ , h versus g_

00

Fluoroelastomer/Cloisite 15A 2.5, 5, 10, 15, 20 phr

Remarks Separation of G into G for breakdown and G for frictional loss

00

G , G versus g_

0

Data Obtained

EPDM/Dellite 67G 5, 10, 20, 30 phr

Materials

TABLE 8.3 (Continued)

[11]

[29]

[28]

[27]

[10]

[9]

References

267

Dependency of rheological properties on mixing time

G0 , G00 versus v

SBS/Cloisite 15A 5 wt.%

SEBS/Cloisite 15A 5 wt.%

(Block copolymers: SBS, four types of SEBS)

Block copolymers/Cloisite 15A 5 wt.%

G0 versus v

G versus strain

0

G0 , |h | versus v



Relative viscosity versus clay loading

Frequency sweep

Temperature sweep

(organoclays: organo-Laponite RD(1), organo-Laponite RD(2), Cloisite 20A)

SEBS/Compatibilizer/Cloisite 20A 2, 5, 7 wt.%

Strain sweep

G0 , tan d at 3.3 Hz, 140 and 117
C

E versus temperature

Effect of preparation method

Effect of preparation method, chemical modification, copolymer microstructure on rheological properties

Effect of compatibilizer on viscosity

Dependency of rheological properties on different organoclays

Dependency of rheological properties on different organoclays

Shear-induced orientation was explained

h/h0 versus l_g 0

Fitting Carreau model to viscosity data to obtain zero-shear viscosity (h0) and relaxation time (l)

h versus g_

TPU/organoclays 1, 3, 5 wt.%

TPU/Cloisite 20A 1, 3, 5, 10 wt.%

TPU/organo-Laponite RD 1, 3, 5, 10 wt. %

TPU/Cloisite 30B 2.5 wt.%

Polyepichlorohydrin/Cloisite 25A 2, 5, 10 wt.%

hE versus g_

Melt strength versus DDR

hE versus strain

[33]

[14]

[32]

[13]

[31]

[30]

[12]

268

RHEOLOGY OF RUBBER–CLAY NANOCOMPOSITES

strain sweep experiments [6,11,14,17,23]. The storage modulus of clay-filled composites was compared with that of carbon black-filled [10,25,26] and silica-filled composites [10,18,25,26]. It was found that the storage modulus of clay composites were higher than that of silica composite, when clay was mixed with polar rubber matrix [18]. The value of the storage modulus of EPDM–organoclay composites at a loading of 10 phr was in between the values of silica and carbon black composites at the same loading [25,26], while the storage modulus of fluoroelastomer composites containing organoclay or pristine clay at a loading of 10 wt.% was lower than that of silica- and carbon black-filled composites at the same loading [10]. The extent of extrudate swell and/or the surface smoothness of extrudate were obtained for NBR [7,22], SBR [7], BR [7], fluoroelastomer [10], BIMS [28] composites. The lower extrudate swell and better surface smoothness were observed upon the addition of organoclay. These are important findings for industrial application of rubber nanocomposites. The significant increase in the storage modulus or steady shear viscosity by adding organoclay was observed for IR [6,17], ENR [18], BR [8,20], NBR [21,22], EPR [23], EPDM [24–26], fluoroelastomer [27], EVA [11,29], polyepichlorohydrin [12], TPU [13,30,31], SEBS [14,32,33] composites. The extent of the increase in the storage modulus or steady shear viscosity was affected by the degree of intercalation and exfoliation of clay that determine the rubber–clay interaction. The improved intercalation and exfoliation of clay are main reasons for a significant increase in the storage modulus and steady shear viscosity. However, the increase in these properties was dependent on the type of organoclay [13,21,27,29–31]. The use of longer alkyl chain as an organic modifier led to the decrease of ML1 þ 4 with the loading of clay in NBR matrix, while the shorter chain did not show such a behavior [21]. No increase in the storage modulus was observed by the addition of some organoclays in fluoroelastomer [27] and EVA [29] matrices. Moreover, a decrease in the storage modulus was observed by the addition of organoclay in TPU at the poor mixing condition [30]. The same behavior was also observed for the TPU composite containing a particular organoclay [13,31]. The presence of compatibilizer in rubber–clay nanocomposites also affects the extent of increase in the storage modulus of EPDM [24] and steady shear viscosity of SEBS [32] composites. The storage modulus of SBS and SEBS composites was also affected by the preparation method, showing the increase by the solution mixing method [14]. In general, an addition of filler into rubber matrix causes an increase in viscosity of the composite. However, the reduction of steady shear viscosity of composites by adding organoclay was observed for NBR [7], SBR [7], and fluoroelastomer [10] composites. This phenomenon is opposite to the result typically observed with conventional fillers, such as carbon black and silica. This behavior was due to the orientation of clay platelets along the flow direction under the high shear rate and the slippage of polymer chains on the clay surface. The shear-induced orientation of organoclay leading to a decrease of the shear viscosity of composites was also observed in hydroxyl-terminated liquid BR [8] and polyepichlorohydrin [12] composites. This unusual rheological behavior seems to appear if the following conditions are satisfied: (1) rheological test is carried out at the high shear rate in a

OVERVIEW OF RHEOLOGICAL THEORIES OF POLYMER–CLAY NANOCOMPOSITES

269

steady-state measurement, instead of a dynamic oscillation [8]; (2) rubber exhibits polar characteristics inducing the orientation of clay platelets, leading to a viscosity reduction [7]; (3) organoclay is used instead of pristine clay, leading to the shearinduced orientation of clay platelets [7]; (4) the organic modifier with a longer chain length decreasing viscosity due to its plasticizing effect [21]. However, these criteria do not always explain the unusual rheological behavior related to the shear-induced orientation of clay, such as the alignment of clay in IR matrix upon imposing preshear leading to an increase or decrease in the storage modulus depending on the concentration of clay [17].

8.4 OVERVIEW OF RHEOLOGICAL THEORIES OF POLYMER–CLAY NANOCOMPOSITES Theoretical studies to predict and understand rheological behavior of polymer composites containing nanoparticles were rarely carried out. White et al. [42] developed shear flow constitutive equations for carbon black-filled rubber and applied this model to one- and two-dimensional shear flow in the die. Isayev and Fan [43] modified Leonov viscoelastic model [44] by introducing stress function caused by a presence of filler and developed a viscoelastic plastic constitutive equation for flow of filled rubber. This model was applied to steady simple shear and transient shear flows for NR and IIR composites filled with CaCO3 and carbon black. A good agreement between experimental and theoretical viscosity was observed. Sobhanie et al. [45] developed a viscoelastic plastic rheological model for particle-filled polymer melts by assuming that the total stress is the sum of stresses in the polymer matrix and the filler network. This theory was compared to the experimental results of PS–CaCO3 composite, showing a good agreement for steady and oscillatory shear flow. Havet and Isayev [46] developed a thermodynamic rheological model of polymer composites containing highly interactive filler by considering stresses caused by flow of polymer and adsorption of polymer on particle surface. This theoretical and experimental data showed a good agreement for PS–silica composite [47]. Kabanemi and Hetu [48] developed a reptation-based rheological model considering polymer–particle interaction and confinement. This model was applied to the rheological behavior of PEO–silica composite, and a good agreement between theory and experimental data was observed. The modeling study of rheology of polymer–clay composites was conducted by Eslami et al. [49]. A mesoscopic rheological model for an incompressible and isothermal suspension of totally exfoliated and homogeneously distributed clay in a polymer melt was developed by reformulating FENE-P model [50]. This model considered polymer–clay and clay–clay interaction. The theoretical prediction showed a fair agreement with the experimental data for PS–PI block copolymer containing organoclay. Eslami et al. [51] modified this model by introducing chain reptation confined into a tube surrounded by polymer and clay to the FENE-P model. The prediction of this model was in a qualitative agreement with the experimental

270

RHEOLOGY OF RUBBER–CLAY NANOCOMPOSITES

data for PS–PI block copolymer containing organoclay. Rajabian et al. [52] predicted rheological behavior in the transient elongational flow of PP–clay composites using the mesoscopic reptation model developed by Eslami and Grmela [53] in combination with a model of ellipsoid particles. The experimental data of PP–clay composite was in a good agreement with the prediction. The theoretical modeling of uniaxially elongational flow of exfoliated PP–clay composites was also studied by Lee et al. [54]. The K-BKZ constitutive equation with two damping functions, so called the WD-FO model [55] and PSM-LT model [56] were used. A good agreement was observed between the prediction by the use of PSM-LT model and the experimental data.

8.5 CONCLUSION AND OUTLOOK The intercalation and exfoliation of clay are the most important phenomena influencing the enhancement of rheological properties in rubber–clay nanocomposites. These phenomena determine the rubber–clay interaction. The papers reviewed here revealed the qualitative correlation between the rheological behavior of composites and the strength of the rubber–clay interaction. The improved rubber–clay interaction as well as intercalation and exfoliation determine the level of viscosity and storage modulus of the composites. The presence of liquid-like or pseudo solid-like behavior is defined by these phenomena. The rheological percolation of clay typically occurs when the transition from the liquid-like to pseudo solid-like behavior is observed. The rheological behavior is also affected by the change in the rubber–clay interaction caused by the polarity of rubber matrix, the use of compatibilizer and the type of clay, as well as the composite preparation method. It was reported that the shear-induced orientation of clay platelets could decrease viscosity of the composite, owing to the slip of rubber chains on the clay platelets aligned along the flow direction. In some cases, the composite containing organoclay even showed the lower viscosity than the unfilled rubber. This phenomenon was found in composites containing a polar rubber matrix and organoclay at particular experimental conditions of flow corresponding to high shear rates. Although the decrease in viscosity of composites by adding clay is unexpected observation, the explanation of this phenomenon was given based on the two competing factors: (1) the rubber–clay interaction leading to the increase in viscosity and (2) the shearinduced orientation of clay leading to the decrease in viscosity. However, this phenomenon is not clearly understood. A brief overview of model equations explaining rheological behavior of nanoparticle-filled polymer composites, including polymer–clay nanocomposites, was presented. It is clear from this overview that theoretical description of their rheological behavior is presently not well understood. Therefore, a significant further development of rheological theories and their extensive comparisons with experimental data on rheology of polymer–clay nanocomposites is required.

REFERENCES

271

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17. Jeon, H. S.; Rameshwaram, J. K.; Kim, G. Structure–property relationships in exfoliated polyisoprene/clay nanocomposites. J. Polym. Sci., Part B: Polym. Phys., 42, 1000 (2004). 18. Vu, Y. T.; Mark, J. E.; Pham, L. H.; Engelhardt, M. Clay nanolayer reinforcement of cis1,4-polyisoprene and epoxidized natural rubber. J. Appl. Polym. Sci., 82, 1391 (2001). 19. Tian, M.; Qu, C.; Feng, Y.; Zhang, L. Structure and properties of fibrillar silicate/SBR composites by direct blend process. J. Mater. Sci., 38, 4917 (2003). 20. Zhu, J.; Wang, X.; Tao, F.; Xue, G.; Chen, T.; Sun, P.; Jin, Q.; Ding, D. Room temperature spontaneous exfoliation of organo-clay in liquid polybutadiene: effect of polymer endgroups and the alkyl tail number of organic modifier. Polymer, 48, 7590 (2007). 21. Kim, J.-T.; Oh, T.-S.; Lee, D.-H. Morphology and rheological properties of nanocomposites based on nitrile rubber and organophilic layered silicates. Polym. Int., 52, 1203 (2003). 22. Ibarra, L.; Rodriguez, A.; Mora, I. Ionic nanocomposites based on XNBR-OMg filled with layered nanoclays. Eur. Polym. J., 43, 753 (2007). 23. Austin, J. R.; Kontopoulou, M. Effect of organoclay content on the rheology, morphology, and physical properties of polyolefin elastomers and their blends with polypropylene. Polym. Eng. Sci., 46, 1491 (2006). 24. Mohammadpour, Y.; Katbab, A. A. Effects of the ethylene–propylene–diene monomer microstructural parameters and interfacial compatibilizer upon the EPDM/montmorillonite nanocomposites microstructure: rheology/permeability correlation. J. Appl. Polym. Sci., 106, 4209 (2007). 25. Tan, H.; Isayev, A. I. Comparative study of silica-, nanoclay- and carbon black-filled ethylene propylene diene monomer (EPDM) rubbers treated by ultrasound. Rubber Chem. Technol., 81, 138 (2008). 26. Tan, H.; Isayev, A. I. Comparative study of silica-, nanoclay- and carbon black-filled EPDM rubbers. J. Appl. Polym. Sci., 109, 767 (2008). 27. Lakshminarayanan, S.; Lin, B.; Gelves, G. A.; Sundararaj, U. Effect of clay surfactant type and clay content on the rheology and morphology of uncured fluoroelastomer/clay nanocomposites prepared by melt-mixing. J. Appl. Polym. Sci., 112, 3597 (2009). 28. Maiti, M.; Sadhu, S.; Bhowmick, A. K. Brominated poly(isobutylene-co-para-methylstyrene) (BIMS)-clay nanocomposites: synthesis and characterization. J. Polym. Sci., Part B: Polym. Phys., 42, 4489 (2004). 29. Riva, A.; Zanetti, M.; Braglia, M.; Camino, G.; Falqui, L. Thermal degradation and rheological behaviour of EVA/montmorillonite nanocomposites. Polym. Degrad. Stab., 77, 299 (2002). 30. Berta, M.; Saiani, A.; Lindsay, C.; Gunaratne, R. Effect of clay dispersion on the rheological properties and flammability of polyurethane–clay nanocomposite elastomers. J. Appl. Polym. Sci., 112, 2847 (2009). 31. Mishra, A. K.; Nando, G. B.; Chattopadhyay, S. Exploring preferential association of Laponite and Cloisite with soft and hard segments in TPU–clay nanocomposite prepared by solution mixing technique. J. Polym. Sci., Part B: Polym. Phys., 46, 2341 (2008). 32. Chen, W.-C.; Lai, S.-M.; Chen, C.-M. Preparation and properties of styrene–ethylene– butylene–styrene block copolymer/clay nanocomposites: I. Effect of clay content and compatibilizer types. Polym. Int., 57, 515 (2008). 33. Carastan, D. J.; Vermogen, A.; Masenelli-Varlot, K.; Demarquette, N. R. Quantification of clay dispersion in nanocomposites of styrenic polymers. Polym. Eng. Sci., 50, 257 (2010).

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34. Thakkar, H.; Goettler, L. A. The effects of DV on the morphology and rheology of TPVs and their nanocomposites. Rubber World, 229, 44 (2003). 35. Mehta, S.; Mirabella, F. M.; Rufener, K.; Bafna, A. Thermoplastic olefin/clay nanocomposites: morphology and mechanical properties. J. Appl. Polym. Sci., 92, 928 (2004). 36. Mousa, A.; Halim, N. A.; Al-Robaidi, A. Rheological and mechanical properties of claythermoplastic elastomers derived from PVC and NBR. Polym.-Plast. Technol. Eng., 45, 513 (2006). 37. Naderi, G.; Lafleur, P. G.; Dubois, C. The influence of matrix viscosity and composition on the morphology, rheology, and mechanical properties of thermoplastic elastomer nanocomposites based on EPDM/PP. Polym. Compos., 29, 1301 (2008). 38. Krieger, I. M. Rheology of monodisperse latexes. Adv. Colloid Interface Sci., 3, 111 (1972). 39. Bicerano, J.; Douglas, J. F.; Brune, D. A. Model for the viscosity of particle dispersions. J. Macromol. Sci., Rev. Macromol. Chem. Phys., C39 561 (1999). 40. Kraus, G. Swelling of filler-reinforced vulcanizates. J. Appl. Polym. Sci., 7, 861 (1963). 41. Sadhu, S.; Bhowmick, A. K. Preparation and properties of nanocomposites based on acrylonitrile–butadiene rubber, styrene–butadiene rubber, and polybutadiene rubber. J. Polym. Sci., Part B: Polym. Phys., 42, 1573 (2004). 42. White, J. L.; Wang, Y.; Isayev, A. I.; Nakajima, N.; Weissert, F. C.; Min, K., Modeling of shear viscosity behavior and extrusion through dies for rubber compounds. Rubber Chem. Technol., 60, 337 (1987). 43. Isayev, A. I.; Fan, X. Viscoelastic plastic constitutive equation for flow of particle filled polymers. J. Rheol., 34, 35 (1990). 44. Leonov, A. I.; Lipkina, E. K.; Paskhin, E. D.; Prokunin, A. N. Theoretical and experimental investigation of shearing in elastic polymer liquids. Rheol. Acta, 15, 411 (1976). 45. Sobhanie, M.; Isayev, A. I.; Fan, Y. Viscoelastic plastic rheological model for particle filled polymer melts. Rheol. Acta, 36, 66 (1997). 46. Havet, G.; Isayev, A. I. A thermodynamic approach to the rheology of highly interactive filler-polymer mixtures: part I—theory. Rheol. Acta, 40, 570 (2001). 47. Havet, G.; Isayev, A. I. A thermodynamic approach to the rheology of highly interactive filler–polymer mixtures. Part II. Comparison with polystyrene/nano silica mixtures. Rheol. Acta, 42, 47 (2003). 48. Kabanemi, K. K.; Hetu, J.-F. A reptation-based model to the dynamics and rheology of linear entangled polymers reinforced with nanoscale rigid particles. J. Non-Newtonian Fluid Mech., 165, 866 (2010). 49. Eslami, H.; Grmela, M.; Bousmina, M. A mesoscopic rheological model of polymer/ layered silicate nanocomposites. J. Rheol., 51, 1189 (2007). 50. Bird, R. B.; Dotson, P. J.; Johnson, N. L. Polymer solution rheology based on a finitely extensible bead-spring chain model. J. Non-Newtonian Fluid Mech., 7, 213 (1980). 51. Eslami, H.; Grmela, M.; Bousmina, M. A mesoscopic tube model of polymer/layered silicate nanocomposites. Rheol. Acta, 48, 317 (2009). 52. Rajabian, M.; Naderi, G.; Dubois, C.; Lafleur, P. G. Measurements and model predictions of transient elongational rheology of polymeric nanocomposites. Rheol. Acta, 49, 105 (2010). 53. Eslami, H.; Grmela, M. Mesoscopic formulation of reptation. Rheol. Acta, 47, 399 (2008).

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54. Lee, S. H.; Kim, S. Y.; Youn, J. R. Rheological behavior and theoretical modeling of uniaxial elongational flow properties of polypropylene/layered silicate nanocomposites. Polym. Compos., 30, 1426 (2009). 55. Feigl, K.; Ottinger, H. C. A numerical study of the flow of a low-density-polyethylene melt in a planar contraction and comparison to experiments. J. Rheol., 40, 21 (1996). 56. Luo, X.-L.; Tanner, R. I. A streamline element scheme for solving viscoelastic flow problems part II: integral constitutive models. J. Non-Newtonian Fluid Mech., 22, 61 (1986).

CHAPTER 9

VULCANIZATION CHARACTERISTICS AND CURING KINETICS OF RUBBER–ORGANOCLAY NANOCOMPOSITES RAQUEL VERDEJO MARIANELLA HERNANDEZ NATACHA BITINIS JOSE MARÍA KENNY MIGUEL ANGEL LOPEZ-MANCHADO

9.1 INTRODUCTION Unvulcanized raw rubbers are entangled high molecular weight viscoelastic liquids, generally not very strong, sticky, easily deformed when warm, brittle when cold, that do not maintain their shape after a large deformation. They are completely soluble in solvents, have a consistency similar to chewing gum and give rise to inelastic deformation being made of long polymeric chains that can move independently to each other. An uncured rubber cannot be used to make articles with a good level of elasticity. The transformation to a useful rubber article such as tires and mechanical goods is due to the discovery of the vulcanization process by Goodyear [1]. In this process, physical and chemical cross-links are formed between the polymer chains giving rise to the formation of a three-dimensional network structure. The long rubber chains with molecular weight of the order of 1
105 g/mol are cross-linked at many points along their length, producing 10–20 cross-links per primary molecule, so that the chains can no longer move independently. As a result, viscoelastic liquids are converted to viscoelastic solids with a high elasticity. They are prone to suffer considerable deformation under stress but upon release of the stress, the rubber article can go back to its original shape, recovering the energy stored during the Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

275

276

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

deformation. In addition, the cross-linked material becomes insoluble in any solvents and it cannot be further processed [2–5]. Rubbers produced by cross-linking amorphous gum elastomers are soft and weak. They require the inclusion of reinforcing fillers, according to a common practice in the rubber industry. The addition of rigid entities can increase stiffness and resistance to fracture due to the appearance of strong interactions between the elastomer and filler surface. The origin of these interactions can be physical or chemical, depending on the chemical nature of the elastomer and on the functionality of the filler surface. Although it is demonstrated that the chemical interaction between fillers like carbon black and elastomer significantly contributes to the reinforcement process, it is not a necessary condition for reinforcement. The main requirement for rubber reinforcement is a sufficiently small particle size, less than 1 mm [6,7]. Small hard domains generate an effective reinforcement, even if the filler–elastomer interaction is weak. As specific surface area increases (i.e., as particulate size decreases), the strength of filled rubber compound usually increases. In composites containing inclusions of 10–100 nm in size, the matrix rubber may show a completely different behavior from bulk gum. Matrix mobility at large strains is reduced because of close proximity to filler surfaces, which is expected to hinder crack initiation and growth. An optimally cross-linked amorphous gum rubber is much stronger, and this is attributed to dissipative mechanisms (internal friction and networks arrangements), which relieve local stress concentration and result in a more uniform distribution of load among network chains. Deformed rubber composites containing inclusions of 10–100 nm in size exhibit much more energy dissipation (hysteresis) owing not only to a high resistance to crack initiation but also to a high difficulty to crack progress. Once a crack begins to grow, it very quickly encounters a rigid filler particle that hinders its progress, inhibiting continued lateral crack growth [8]. Energy dissipation plays an essential role in this mechanism of reinforcement. Hysteresis in filled rubber has been attributed primarily to slippage of adsorbed chains over filler surfaces, but may also have a contribution from the slippage of unabsorbed chains, which are entangled with adsorbed chains. In both molecular mechanisms, the extent of hysteresis increases with increased filler surface area, that is, with smaller particles. So, the essential importance of particle size is now apparent.

9.2 VULCANIZATION REACTION The vulcanization reaction is an irreversible process that gives rise to the formation of a three-dimensional network structure of polymer chains chemically bonded to one another, converting a plastic material into an elastic one. This transformation can occur through various chemical mechanisms such as nucleophilic substitution, end linking, addition or condensation, free radical coupling and ring opening reactions. The physical properties of a rubber material are largely depending on the state of cure of the final compound. Thus, a good understanding of the vulcanization process is necessary to obtain useful rubber articles.

VULCANIZATION REACTION

277

Increase Equilibrium

Modulus

Reversion

Induction

Curing

Overcure

FIGURE 9.1 Typical curing curve of a rubber. Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

The cure process consists essentially in the heat transfer by conducting through the rubber mass and the cure reaction which starts around 120–170 C depending on the nature of the elastomer. If the mould temperature is too low, the reaction is not properly achieved, and the final material exhibits little interest for industrial application. On the other hand, if the temperature is too high, the final material is burnt, loosing the quality of the elasticity. Thus, it is necessary to obtain a good knowledge on the kinetic parameters of the cure reaction. The reader interested in this subject is referred to reviews on vulcanization and cross-linking in elastomers [9–18]. The vulcanization curves of a rubber compound are currently obtained by monitoring the increase of the torque required to maintain a given amplitude of oscillation at a given temperature in an oscillating disc rheometer (ODR). It has been assumed that the increase in torque during vulcanization is proportional to the number of cross-links formed per unit of volume of rubber. The torque is automatically plotted versus time to give a so-called cure curve or rheograph. A typical curing curve of a rubber is shown in Figure 9.1. The curve gives a rather complete picture of the overall kinetics of cross-link formation and it is therefore extensively used to control the quality and uniformity of rubber stocks. Three regions are clearly observed: the first region is the scorch delay or induction period, during which most of the accelerator reactions occur. The second period is due to the curing reaction, during which the network structure is formed. It defines the curing speed, that is, the greater or lesser degree of activation of the reaction between the rubber and the cross-linking system. In the last period, once the optimum vulcanization is achieved, a plateau is observed. The network matures by overcuring reversion, equilibrium, or additional but slower cross-linking, depending on the nature of the compound. The overcuring reversion corresponds to a loss of the

278

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

mechanical behavior of the material due to the breakdown of bonds, in particular, polysulfide bonds. The main parameters obtained from the ODR curve provide information about the reaction kinetics and the cross-link density and are detailed below: .

.

.

Scorch Time (tS2). It is the time required for the torque value to increase by 2 in. pounds over the minimum. It indicates the time available before onset of vulcanization and, it provides a good assessment of the scorch safety of a rubber compound. Lower scorch time values correlate with an increased likelihood of premature cross-linking (precocity). Optimum Cure Time (t90). It is the time in which 90% of the delta torque is reached. This is a useful estimate of the overall cure rate at a given temperature. As overheating the material could lead to reversion processes. Cure Rate Index (CRI). It is a direct measure of the fast curing nature of rubber compounds and can be calculated using the following expression CRI ¼ 100=ðt90  tS2 Þ

.

. .

ð9:1Þ

The Minimum Torque Value (S0 min ). It represents an index of material viscosity and can be related to the dispersion of the filler and the filler–polymer interaction. In general, it indicates the extent of the mastication. The Maximum Torque (S0 min ). It represents the highest level of cross-linking possible at a given vulcanization temperature. The Delta Torque (DS ¼ S0 max  S0 min ). It is the difference between the maximum and the minimum torque, and it is a quantitative assessment of the crosslinking density of a vulcanizate. So, it is an efficient means of measuring the effects of additives on the cure efficiency.

9.3 RUBBER CROSS-LINKING SYSTEMS Several cross-linking systems such as sulfur, sulfur donors, organic peroxide, phenolic resins, benzenoquinone, bismaleimides, amino acids, azide compounds, aldehydes, silane, electron-beams, UV radiation, g rays, ultrasonic waves, microwave, and so on, have been used for the rubber vulcanization. Among these, sulfurand peroxides-based systems are the most widely used in the rubber industry. 9.3.1 Sulfur Vulcanization The traditional yet most common vulcanization method for polydienes, including natural rubber (NR), is the sulfur-based one [19,20]. The rubber reaction with only sulfur is very slow and has a very poor efficiency and, as a consequence, it is inadmissible at industrial level. Sulfur vulcanization compound is generally used alongside with a wide range of specific chemicals that provide the desired course of

RUBBER CROSS-LINKING SYSTEMS

279

TABLE 9.1 Classification of the Accelerators as a Function of Their Relative Curing Speeds [9] Type

Curing Speed

Guanidines (DPG, DOTG) Thiazoles (MBT, MBTS, ZMBT) Thiurams (TMTD, TMTM, TETD, DPTTS) Dithiocarbamates, xanthates (ZDBC, NDBC)

Slow Moderate Fast Very fast

vulcanization and target product properties. A standard accelerated sulfur vulcanization formulation includes the sulfur as cross-linking agent, activators, one or more accelerators, retardants and prevulcanization inhibitors. The addition of activators (usually zinc oxide and stearic acid) and accelerants (thiurams, sulfenamides, guanidines, etc.) offers many advantages for the curing process, such as shortening of time (periods as short as 2–5 min) and lowering of temperatures, thus reducing thermal and oxidative degradation. Accelerators allow a more efficient reaction with a perfect control over the process and, in addition, the reduction of the amount of sulfur used, avoiding the bloom of the unreacted one, finally improving the compound ageing. Both scorch resistance and cross-linking rate are affected by the type of accelerator. So, special attention has to be taken into account to choose the proper accelerators to obtain appropriate rubber network. Table 9.1 illustrates the most important classes of accelerators used in the rubber industry and their relative curing speeds [9]. Frequently, combinations of accelerators are used to an accelerator, named primary, a small amount is added of another accelerator, called secondary. A practical example is the blocking of the active group of 2-mercaptobenzothiazole (MBT) in order to reduce its precocity, by the addition of an amine. Thus, at a vulcanization temperature, a synergistic effect MBT/amine is produced. The mechanism of accelerated sulfur vulcanization is still controversial partly due to the complexity of the formulation; in particular, the mechanisms of individual reactions are still not completed known or are in dispute. One of the fundamental issues is whether the main reaction pathway is via ionic or free-radical mechanism (see Schemes 9.1 and 9.2, respectively). However, the mechanism of accelerated sulfur vulcanization can be generalized as follows: [9,12] the accelerator reacts with the sulfur to give a more reactive polysulfide compound. At this point, the available Zn2þ (formed by the interaction with the stearic acid) forms complexes with the accelerators and the polysulfide compounds. Finally, the rubber polysulfides react—either directly or through an intermediate—to give sulfur cross-links, of the cyclic or C–Sx–C type [14]. Different types of structural groups have been proposed for vulcanized networks. These include cyclic sulfide units, pendant side groups, isomerized double bonds, and conjugated unsaturated bonds, in addition to a variety of sulfur cross-linking units. The accelerated sulfur vulcanization systems can be classified into three groups, as a function of the sulfur level in the formulation and of the accelerator/sulfur ratio (Table 9.2): conventional vulcanization, semiefficient (semi-EV), and

280

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

Initiation

RSa+ + RSb–

Polysulfide (RS aSbR) CH2

Propagation

RSa+

+

CH2 C C

6 +

CH2

C C

RSa +

C C 6

CH proton

C

transfer

RSa C

CH2 C+ C H

+

Crosslink CH C +

CH2 hydride transfer

C H RSa C

+ C 8

Crosslink CH2 C S8 + ( C H

7 + S8

RSa +)

CH C C S8 + ( RSa +)

8 + S8

Termination

}

7 8

+

}

RSa

+

}

+

RSb–

nonchain carriers

SCHEME 9.1 Ionic mechanism of sulfur vulcanization. Reprinted with permission from Ref. [9]. Copyright 1997 Elsevier.

TABLE 9.2 Classification of the Sulfur Vulcanization Systems [19] Type Conventional Semi-EV EV

Sulfur (phr)

Accelerator (phr)

A/S Ratio

2.0–3.5 1.0–1.7 0.4–0.8

1.2–0.4 2.5–1.2 5.0–2.0

0.1–0.6 0.7–2.5 2.5–12

phr: parts per hundred of rubber.

RUBBER CROSS-LINKING SYSTEMS

homolytic (S)8

heterolytic

•S

(S)6 S•

+S

(S)6 S–

(1)

R S2X R

(2)

S8+2 olefin R SX

C

• ) 2 R SX

SX R (

+ RSX •

(3)

C SX R

C

281

C•

(4)

4

4

C S x-y R

RSX•

C (S)y•

4 + S6

C SX R

RSX •

C (S)6 •

4 + R S8 H

C SX

(5)

(6)

(7)

R + RSX •

C H Cross link

C C C H + RSX• (

C C C • + RSXH

)

•C

(8)

C C 5

5

C C C S6 •

+ S6

RSX •

(9)

H C C C

RSXH

H C H C

C SX

(10) R

SCHEME 9.2 Free-radical mechanism of sulfur vulcanization. Reprinted with permission from Ref. [9]. Copyright 1997 Elsevier.

282

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

efficient (EV) sulfur cure system. The conventional systems have a relatively high sulfur to accelerators (2.5 w/w), and would result in predominantly polysulfide bridges at optimum cure [21], whereas an EV system is characterized by a very low sulfur to accelerator ratio (0.15 w/w), giving higher cross-link density because the cross-links are shorter. They form networks containing mainly monosulfidic crosslinks and fewer polysulfides ones, and exhibit a low degree of main chain modifications. These EV systems suffer with respect to fatigue resistance but exhibit an excellent reversion resistance and good resistance to thermal and oxidative aging. 9.3.2 Peroxide Vulcanization Sulfur vulcanization requires the presence of double bonds on the elastomer backbone. For saturated rubbers such as ethylene–propylene rubber (EPR) and silicones, it is necessary to use other cross-linking agents, such as organic peroxides. A wide variety of peroxides can be used to cross-link most types of elastomers [10,12,22]. The peroxides are able to create carbon–carbon cross-links via free radicals even on overcure [22,23]. The peroxide efficiency is a term that denotes the number of moles of cross-links that can be formed from a mole of peroxide. The driving force of peroxide vulcanization is free radicals. The peroxide rubber vulcanization reaction consists of three steps as shown in Scheme 9.3: (i) The process is initiated by the thermal homolytic scission of a peroxide molecule to form two highly energy alkoxy radicals [10]. This is a first-order reaction. (ii) These radicals react with the polymer chain removing the more labile allylic hydrogen atoms from the polymer (hydrogen heat of light

ROOR

2RO •

RO • + P H

ROH + P •

2P •

P P Crosslink

2P •

P(–H) + P (+H) R

RO• + ∼∼ CH2

C CH ∼∼

R • ROH + ∼∼ CH C CH∼∼

R • ∼∼ CH C CH ∼∼ • CH2

CH3 RO • +

CH2

C CH

CH2

C CH ∼∼

CH2

∼∼ CH2

•

C CH ∼∼

SCHEME 9.3 Peroxide cross-linking mechanism. Reprinted with permission from Ref. [9]. Copyright 1997 Elsevier.

THE ROLE OF ORGANOCLAY ON VULCANIZATION REACTION

283

abstraction). (iii) Then, two radicals on adjacent polymer chains couple to form a cross-link (radical coupling). The cross-link is a carbon–carbon bond between the chains. This bond is strong and rigid compared to the sulfur cross-link. The sulfur cross-link is less thermally stable than the carbon–carbon bond. Neither the peroxide nor by-products from the vulcanization process are part of the crosslinks. Hence, the inherent stability of the polymer is not reduced by cross-linking with peroxides [24–26]. In opposition to the sulfur system, the vulcanization rate with organic peroxide cannot be modified by the addition of other substances, because it only depends on the peroxide half-time at a given temperature. The half-time is the time required for half of the peroxide to decompose at the reaction temperature. Therefore, the half-time is related to the cure temperature. In general, for typical peroxides, the half-time drops to about one-third of its value for each 10 C increase of the temperature [27]. Other disadvantages of radical vulcanization are the secondary reactions of polymer scission, the beta cleavage of the oxy radicals, and the destruction of macroradicals by molecular disproportionation. These reactions compete with themselves and consume radicals in nonproductive way, reducing the efficiency of the vulcanization process and producing unwanted effects.

9.4 THE ROLE OF ORGANOCLAY ON VULCANIZATION REACTION Hybrid organic–inorganic nanocomposites consisting of a polymeric matrix and a layered silicate have inspired scientists to a range of potential applications. Due to their nanometer phase dimensions, nanocomposites exhibit significant improvements in physical and mechanical properties in relation to the polymer host. The addition of just a low percentage of nanolayered inorganic particles can increase the stiffness and strength with a minimal loss in ductility and impact resistance, decrease the permeability and swelling in solvents, improve the abrasion, flame resistance and thermal endurance, with an enhancement in electrical conductivity and optical properties [29–32]. In the particular case of elastomers, and in order to achieve a material with desired properties, it is necessary to understand the effect of nanoparticles in the vulcanization reaction. As already mentioned, the vulcanization reaction is a critical step in obtaining useful rubber articles. Several researchers [33–70] have evaluated the influence of organoclay on the rubber vulcanization characteristics. In general, it has been widely reported in the literature that, regardless of the nature and polarity of the elastomer, the organoclay exerts two effects on the rubber cure reaction. First, it behaves as a catalyst, accelerating the vulcanization reaction since vulcanization times, scorch time and optimum cure time are simultaneously and sensibly decreased. A low scorch safety may represent a processing problem but a high cure speed, in general, can be considered beneficial [69]. Second, it increases the delta torque value indicating a strong influence of organoclay on the cross-linking density of the rubber compound. These effects are observed even at low organoclay concentrations in the composite

284

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

(1 wt.%), significantly lower than those used with conventional fillers, such as carbon black or silica. In addition, Zhang et al. [67] demonstrated an improvement of the antireversion of natural rubber/chloroprene (NR/CR) blends by addition of organoclay. The authors attributed this behavior to the fact that the organoclay stays intercalated mainly in the CR phase, improving the thermal stability of the blend and giving a barrier wall for the interdiffusion of both oxygen and hydrochloride. The accelerant effect of the organoclay is essentially attributed to the presence of the amine modifier inserted into the clay galleries. It is well known that the amine groups activate functional groups of the accelerants, for instance benzothiazyl disulfide, giving rise to a synergic effect that leads to a faster and more effective vulcanization reaction. However, it is worth mentioning that this effect was still more noticeable when the octadecylamine was intercalated between the silicate galleries [37]. Although a complete knowledge of the exact nature of the interaction mechanism of the octadecylamine in the vulcanization system is still lacking, several authors [41,46,50,53] have proposed the formation of a zinc coordination complex during curing in which the amine groups of the modifier intercalated in the clay gallery and the sulfur participate. The existence of a Zn–sulfur–amine complex in rubber vulcanization systems has been suggested in the previous literature [71–73]. Wang et al. [70] proposed the formation of key accelerator species, shown in Figure 9.2a, during the induction period for the sulfur/sulfenamides vulcanization system [18]. This activated intermediate can react with an amine to form the chemical compound represented in Figure 9.2b [70], where the ligand denotes the amine. This chelate is more active than the sulfenamides accelerator. In general, it has been observed that rubber chains partially penetrate into the silicate interlayer space during compounding in an internal mixer or roll mill, before the vulcanization reaction. This is reflected in a slightly increase of the basal spacing of clay as detected by X-ray. So, for instance, Usuki et al. [35] observed that the N

N (a)

C Sx Zn Sy C S

S Ligand

N

N (b)

C Sx Zn Sy C S

S Ligand

FIGURE 9.2 Key accelerator species formed during the vulcanization induction period for a sulfur/sulfenamides vulcanization system (a) and reaction product with an amine (b) (see Ref. [70]).

THE ROLE OF ORGANOCLAY ON VULCANIZATION REACTION

285

intercalation of rubber molecules (in this case, ethylene–propylene–diene terpolymer rubber, EPDM) into the organophilic clay occurred during mixing in a roll mill, while the complete exfoliation process succeeded during vulcanization process. The conceptual image of EPDM–clay hybrid material prepared via vulcanization process is illustrated in Figure 9.3. Furthermore, the authors observed that the nanocomposite morphology was dependent on the accelerator used in the rubber recipe. EPDM–clay hybrids were successfully prepared when thiurams and dithiocarbamates were used as vulcanization accelerators, while the dispersibility of clay was insufficient in presence of thiourea, thiazole, and sulfenamide. The authors explained this fact as due to thiuram and dithiocarbamate accelerators dissociation into radicals at high temperatures. These radicals link with carbon atoms in EPDM backbone to polarize EPDM molecules that intercalate into the clay galleries through hydrogen bonds between the polar EPDM and the clay surface (Figure 9.3). On the other hand, Zheng et al. [74] and Gatos and Karger-Kocsis [49] observed that the addition of the curative agents in a second stage of compounding led to further clay layers separation. As shown in Figure 9.4, trace c, an increase of 0.76 nm in the interlayer distance is observed after the addition of curatives to an organoclay/ EPDM system in an open roll mill [49]. This phenomenon indicates a prolongation of the compounding. The authors concluded that the expansion of the organoclay might be related to the adsorption of the curatives in the clay interlayer space. Moreover, results seem to indicate that the curatives intercalated between two opposite layers are involved in the vulcanization reaction, favoring the intercalation/ exfoliation phenomena. Vulcanization curatives such as zinc oxide (ZnO), accelerant (MBTS), stearic acid, and sulfur are as well polar low molecular mass compounds that may easily penetrate between organoclay layers resulting in the rubber cross-linking inside the layer galleries. It is important to note that Gatos et al. [46], analyzing sulfur-cured HNBR– organoclay nanocomposites, observed that a primary amine, used as montmorillonite modifier, reacted with the sulfur curatives during vulcanization and formed a confined/deintercalated structure in addition to an intercalated/exfoliated one. However, when a quaternary ammonium intercalant was used, no deintercalation was detected. For it, the authors concluded that, in the Zn–sulfur–amine complex formation, only primary amines are involved. However, the opposite effect has been reported by other authors for similar rubber–organoclay nanocomposites. That is, the addition of the organoclay retarded the curing process, increasing the scorch time and optimum cure time [74–79]. So, Chang et al. [75] observed that the vulcanization time (t90) for EPDM– organo-montmorillonite hybrid nanocomposites was prolonged with the progressive addition of the organoclay in the compound, presumable due to the absorption of the curing agents on the filler surface. Hwang and Wei [78] reported that the optimum cure time of NBR–organoclay nanocomposites sensibly increased with the filler content in the nanocomposite and attributed this fact to the acidity of the organosilicate, which exhausts some decomposed accelerator free radicals in the crosslinking reaction, retarding the curing reaction.

286

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

FIGURE 9.3 Conceptual figure of vulcanization process of EPDM and intercalation into clay gallery. Reprinted with permission from Ref. [35]. Copyright 2002 Elsevier.

THE ROLE OF ORGANOCLAY ON VULCANIZATION REACTION

287

1.53 nm

3.30 nm

Intensity [a.u.]

(d)

3.08 nm (c) 2.32 nm (b)

2.10 nm

(a) 0

2

6

4

8

10

2θ (º)

FIGURE 9.4 XRD patterns of (a) the MMT-PRIM (primary amine) and (b–d) the EPDM/ MMT-PRIM (10 phr) nanocomposite at the different stages of processing: (a) MMT-PRIM powder, (b) mixing of the EPDM and the MMT-PRIM (10 phr) in the internal mixer, (c) after addition of the curatives on the open mill, and (d) after vulcanization at 160 C. Reprinted with permission from Ref. [49]. Copyright 2004 Elsevier.

A broad consensus in the scientific community exists regarding the reinforcing effect of organoclay in elastomeric matrices. It is assumed that the addition of inorganic nanoparticles decreases the minimum torque due to the amine acting as a lubricant for rubber compounds, and significantly increasing the torque value, which is intimately related to a higher cross-linking density of the system due to the formation of new filler–rubber interactions. This behavior can be easily corroborated by equilibrium swelling measurements in suitable solvents by application of the Flory–Rehner equation [80]:     f  lnð1  fr Þ þ fr þ xf2r ¼ V0 n fr1=3  r ð9:2Þ 2 where fr is the volume fraction of polymer in the swollen mass, V0 is the molar volume of the solvent, and x is the Flory–Huggins polymer interaction term. Most of these works have been focused on conventional accelerated-sulfur cure systems. The vulcanizing characteristics of rubber–organoclay nanocomposites cured with peroxide as cross-linking agent have been less studied [81–84]. Nah et al. [81] analyzed NBR–organoclay nanocomposites and did not find differences neither on curing times nor on the value of maximum torque with the addition of organoclay. The authors attributed this behavior to the strong reactivity of peroxidetype cure agent with the double bonds in NBR backbone. Similar results were reported by Zheng et al. [83] when analyzing EPDM–montmorillonite nanocomposites. The authors observed that the optimum cure time of the EPDM composites

288

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

cured with the peroxide system changed very little with increasing OMMT content, and it was nearly the same as that of gum EPDM vulcanizate. On the other hand, the torque value of the EPDM–OMMT composite was independent of the OMMT content, indicating that the organoclay had little effect on the cross-linking density of the EPDM composites. However, Das et al. [84] found for a similar NBR–organoclay hybrid nanocomposite that the addition of 5 phr of organoclay increased the torque delta and reduced the optimum cure time. Also, they observed that these effects were remarkably higher in the case of a peroxide-cured system as compared with a typical sulfur-cured system. An explication to this fact is omitted but it is of interest to note that the authors observed a better intercalation of the NBR chains when the rubber was crosslinked by sulfur curing as compared to peroxide-cured system. In addition, the authors revealed that the stearic acid had a marked effect on the intercalation process of the organoclay in the rubber matrix when the sulfur-curing system was used for the vulcanization of NBR. The stearic acid is a small molecule that can easily penetrate in the silicate gallery and promote the dispersion of the clay. Furthermore, the formation of chemical bonds between the acid groups of stearic acid with the silanol groups of the silicate was expected. These interactions favored the intercalation of the polymer chains into the silicate galleries. 9.4.1 Influence of Organoclay Structural Characteristics on Rubber Vulcanization The effect of the structural characteristics of layered silicates on the rubber vulcanization reaction has not been systematically studied. Besides the filler percentage in the composite, key aspects with a direct influence on rubber vulcanization reaction and hence on the properties of the finished material appear to be structural organoclay parameters, such as (i) type of organic modifier used as intercalant: ammonium or phosphonium salts; (ii) structure of the organic modifier: primary or quaternary salts; (iii) length of the modifier chain (number of carbon atoms); (iv) cationic exchange capacity of the silicate, CEC or concentration of organic modifier. Most significant findings reported in the literature regarding the effect of the structural characteristics of layered silicates on the rubber vulcanization behavior are briefly discussed. Kim et al. [54] analyzed the effect of the ammonium modifier concentration on the vulcanization behavior of polybutadiene rubber–organoclay nanocomposites. They observed that the organoclay with a higher cation exchange capacity (CEC) accommodated higher amount of intercalant in the clay galleries and promoted the vulcanization reaction. As a result, larger torque differences and faster

THE ROLE OF ORGANOCLAY ON VULCANIZATION REACTION

289

vulcanization rate for the nanocomposite were observed. On the other hand, at the same modifier concentration in the intercalant, the organoclay with higher hydrophobicity increased the torque value and accelerated the vulcanization reaction. Avalos et al. [85] studied the influence of two quaternary phosphonium salts (aromatic and aliphatic) on the vulcanization kinetic of natural rubber–clay nanocomposites. They observed that due to its lower rigidity, the aliphatic salt was easier to intercalate into the clay galleries than aromatic salt giving rise to a higher interlayer distance, which facilitated the rubber intercalation into the silicate galleries. Therefore, a nanocomposite with an exfoliated structure was obtained. The aliphatic phosphonium salt modified montmorillonite behaved as accelerant agent decreasing the cure time and also increasing the torque value. As a result, a higher cross-linking degree was observed in the nanocomposite, which gave rise to a sensible improvement of the processing and physical characteristics of the material. However, no significant changes on the natural rubber vulcanization behavior in the presence of aromatic quaternary phosphonium salts were observed. The effect of the amine intercalant type, primary or quaternary ammonium salt on the cure characteristics of NR–organoclay nanocomposites was also analyzed by Magaraphan et al. [36]. Scorch times and optimum cure times were significantly reduced compared to pure NR and unmodified Na-MMT/NR counterparts, thereby leading to an increase in the cure rates. However, the effects of primary and quaternary ammonium salt modifying agents on cure characteristics were different. The authors observed that the nanocomposites prepared with long quaternary ammonium salts showed faster cure times in relation to those containing primary alkyl amines. The same trend was observed in the reduction of scorch times. Moreover, the authors analyzed the effect of the length of hydrocarbon in the alkyl amine on the properties of NR–organoclay nanocomposite. They demonstrated that the longer organic modifier resulted in a better expansion of the interlayer spacing giving rise to an exfoliated structure and thus enhanced mechanical and physical properties of the nanocomposite. However, no significant differences on rubber cure characteristics were observed. Another important aspect to be analyzed is the effect of the organoclay content in the nanocomposite. It is hard to define an optimal filler concentration, since it is dependent on the type of elastomer, vulcanization system and degree of filler dispersion in the polymer. In general, it can be deduced that the nucleating effect of organoclay is reached by concentrations of up to 5 wt.%. Further addition of the organoclay does not affect or delay the vulcanization process, probably due to a partial absorption of curing agents on the filler surface. However, the maximum reinforcing effect of organoclay, deduced as an increase of torque value is attained at higher levels of organoclay in the nanocomposite, approximately 7 wt.%. As the organoclay concentration in the nanocomposite is increased a slight decrease in the torque value is obtained due to the formation of agglomerates. However, a more detailed study is required to understand the role of the organoclay in the rubber vulcanization reaction.

290

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

9.5 VULCANIZATION KINETICS OF RUBBER–ORGANOCLAY NANOCOMPOSITES For many years, although the cross-linking process and the use of fillers were known to be necessary to obtain good performance properties, the vulcanization reaction mechanism remained almost unstudied. However, the need to improve rubber materials led to pay special attention to the role of the filler in vulcanization as well as to determine the kinetic parameters that have a major influence on this process. Several models have been applied to study the kinetics of rubber vulcanization. These can be classified into two groups: phenomenological approaches [86–88] that usually successfully fit the cure curve, but do not provide information about the chemistry occurred during the reaction, and the mechanistic approaches [89–93] whose main objective is to understand the chemical reaction and mechanisms that take place during the vulcanization reaction. However, in most of these models, the filler is omitted from the formulation. Numerous analytical techniques such as nuclear magnetic resonance (NMR) [94,95], mass spectrometry (MS) [96], infrared (IR) [97,98], and Raman spectroscopy [99–101] have been applied, but always with difficulties, and they cannot discriminate between different reactions that take place at the same time. Popular techniques used to study the kinetics of rubber vulcanization include differential scanning calorimetry (DSC) and ODR [89,102,103]. To date, the kinetics of vulcanization of rubber nanocomposites has been scantily studied [37,43,47,59]. The reported studies have been focused on analyzing the effect of layered silicates on the kinetic parameters and activation energy of curing reaction. The cure kinetic studies on natural rubber–organoclay nanocomposites [37], fluoroelastomer–clay nanocomposites [43], nitrile rubber–layered clay nanocomposites [47], and acrylate rubber–organoclay nanocomposites [59] showed the suitability of the autocatalytic model for analyzing the cure parameters of rubber– clay nanocomposites. Lopez-Manchado et al. [37] analyzed the effect of the incorporation of a layered silicate on the vulcanization kinetics of natural rubber in an ODR and DSC under dynamic and isothermal conditions. DSC is a major achievement in the field of analytical tools for the determination of the energy required during vulcanization reaction. Analysis by DSC assumes that the heat of the reaction is only due to a single curing reaction and is proportional to the extent of the reaction. The rheometer measurements are based on the fact that the crosslinking density is proportional to the stiffness of the rubber. Therefore, the degree of curing (a) can be calculated from the heat-flow peak of a DSC curve or from the dynamic viscoelastic properties obtained in curemeter testing [90]. So, a can be easily calculated from the following equations. For DSC a ¼ DHt =DH1

ð9:3Þ

VULCANIZATION KINETICS OF RUBBER–ORGANOCLAY NANOCOMPOSITES

291

For the curemeter test a ¼ ðFt  F0 =F1  F0 Þ

ð9:4Þ

where DHt is the accumulated heat evolved by time t, DH1 is the total amount of heat generated during the reaction, and F0, Ft, and F1 are the torque values at time zero, at curing time t, and at the end of the vulcanization process. It is assumed that the equation describing the thermal behavior of an elastomer during the vulcanization process is obtained by the appraisal at each instant of a thermal balance, where the diffusion of heat and source terms originated from the exothermic reaction occurring during the vulcanization phenomenon are taken into account. The dynamics of the reaction, after the induction time (ti) has elapsed, are modeled by means of a differential equation with regard to time, where the influence of the temperature and vulcanization rate are separated, as in many other classical chemical reactions. Thus, the equation governing the vulcanization rate may be written as follows: da ¼ KðTÞf ðaÞ dt

ð9:5Þ

where t is the time, T is the temperature, and K is the kinetic constant. The function K (T) is described by the Arrhenius expression: KðTÞ ¼ K0

expð Ea =RTÞ

ð9:6Þ

where K0 is the preexponential factor, Ea is the activation energy, and R is the universal gas constant. From Eqs. (9.3) and (9.4), the following relationship can be deduced: da ¼ K0 f ðaÞ expð Ea =RTÞ dt

ð9:7Þ

The function f(a) depends on curing reaction mechanisms. So, in a single reaction, the function f(a) assumes that the rate of conversion is proportional to the fraction of noncross-linked material [104]. f ðaÞ ¼ ð1  aÞn ;

n1

ð9:8Þ

where n is the reaction order. A more complex form of the kinetic equation assumes the so-called autocatalytic model, which is given by [83] f ðaÞ ¼ am ð1  aÞn ; where m also denotes the reaction order.

0  m  1;

n1

ð9:9Þ

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VULCANIZATION CHARACTERISTICS AND CURING KINETIC

In an n-order reaction, the maximum rate of cure occurs at the beginning, when the concentration of reactive groups is maximum, whereas in an autocatalytic reaction, the maximum rate of cure is given for a conversion degree other than zero because the reaction is promoted by the same products of the reaction [105]. The curing reaction of natural rubber was better fitted by an autocatalytic reaction so, to analyze the curing reaction of NR in the presence of nanostructured silicates, the following expression was proposed: da ¼ KðtÞam ð1  aÞn dt

ð9:10Þ

where n and m are the reaction orders. The choice of this specific model was based on the experimental results obtained in the isothermal tests, which confirmed that the reaction rate was equal to zero at both ends of the process (a ¼ 0 and 1) as predicted by Eq. (9.8). The curing characteristics and the kinetic parameters obtained from curemeter testing for NR–organoclay nanocomposites were summarized in Table 9.3. First, the authors observed that both, cure times, scorch time and optimum cure time decreased with increasing cure temperature. An increase of 10 C in the temperature involved a decrease of almost two times in the optimum cure time. As expected, the torque value gradually decreased as the cure temperature increased. On the other hand, to evaluate the effect of the incorporation of the clay (unmodified and organically modified) on

TABLE 9.3 Kinetic Parameters Obtained from Curemeter Testing at Several Cure Temperatures for NR–Organoclay Nanocomposites Material NR

NR–unmodified clay

NR–octadecylamine

NR–organoclay

Tc ( C)

ti (min)

t97 (min)

DS (dN m)

K

n

m

160 170 180 190 160 170 180 190 160 170 180 190 160 170 180 190

3.25 1.75 1.00 0.65 2.35 1.30 0.75 0.55 1.50 1.00 0.60 0.50 0.80 0.65 0.45 0.35

9.48 4.88 2.57 1.49 8.58 4.24 2.10 1.27 3.64 2.12 1.32 0.90 2.34 1.42 0.95 0.78

4.30 4.12 3.80 3.64 3.40 3.18 2.83 2.66 5.78 5.59 5.36 5.03 8.71 8.36 7.65 6.85

0.24 0.56 0.93 1.37 0.28 0.63 1.07 1.45 0.58 1.12 1.88 2.34 0.71 1.22 2.03 2.47

0.68 0.51 0.58 0.52 0.53 0.44 0.46 0.55 0.58 0.48 0.52 0.53 0.62 0.44 0.51 0.55

1.54 1.38 1.49 1.29 1.60 1.54 1.52 1.49 1.57 1.48 1.46 1.55 1.55 1.45 1.44 1.53

Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

Ea (kJ/mol)

95.9

91.4

78.7

71.1

VULCANIZATION KINETICS OF RUBBER–ORGANOCLAY NANOCOMPOSITES

10

293

NR NR–clay NR–octadecylamine NR–organoclay

S′ (dN m)

8

6

4

2

0

0

5

10

15

20

25

30

Time (min)

FIGURE 9.5 Influence of the clay on the NR rheometer curves at 160 C. Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

the cure characteristics of NR, the vulcanization curves at 160 C of all the studied materials were represented in Figure 9.5. It was deduced that the unmodified clay slightly decreased the cure time of pure NR. However, both times were noticeably reduced when the organoclay was added into the elastomer. Similar behavior was observed at all the tested temperatures, reflecting that the organoclay behaved as a vulcanizing agent for NR, leading to a noticeable increase in the vulcanization reaction. As already mentioned in this chapter, this nucleating effect is mainly attributed to the presence of amine groups in the clay galleries. To corroborate this premise, the authors evaluated the effect of the octadecylamine in the absence of the clay. As was observed, the octadecylamine sensibly reduced the vulcanization times and slightly increased the torque value. It is known that the combination of a benzothiazyl accelerant with an amine produces a particular accelerating effect on the rubber vulcanization reaction. Nevertheless, it is worth noting that a further prominent effect on NR curing was observed in the presence of organoclay. In fact, the intercalation of the octadecylamine within the silicate gallery facilitated the vulcanization reaction, which showed a noticeable decrease in the required time for NR vulcanization compared to the blend with only octadecylamine. On the other hand, the authors demonstrated that the incorporation of organoclay gave rise to a sensible increase in the torque delta in relation to pure NR and unmodified clay–NR composite, indicating the formation of a higher cross-links number. This reinforcing effect was attributed to a strong filler–rubber interaction due to the larger intercalation of rubber chains into the silicate galleries. The kinetic parameters of the curing reaction of NR and its composites obtained from a rheometer were determined by application of Eq. (9.8). The values of K, n, and m at the different tested temperatures were calculated through linear multiple

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

294

2.5 y = m1*(m0/m2)*((1-m0)*m3) Value Error 1.2237 0.63707 m1 m2 0.44388 0.046952 m3 1.4545 0.11569 NA Chisq 0.031272 NA R 0.99241

dα /dt

2 1.5 1 0.5 0

0

0.2

0.6

0.4

0.8

1

α

FIGURE 9.6 Derivative of a as a function of a for NR–organoclay nanocomposites at 170 C: (*) experimental results and (—) resulted predicted from Eq. (9.10). Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

regression analysis of the experimental data, as shown in Figure 9.6. A plot of ln K versus 1/T (Figure 9.7) gives a straight line of slope  Ea =RT, where Ea is the activation energy generated during the vulcanization reaction and R is the gas universal constant. The complete set of kinetic parameters were reported in Table 9.3. The values of K were in concordance with those of optimum cure time. That is, the organoclay behaves as an effective vulcanizing agent for NR vulcanization, showing a noticeable increase in K at all the tested temperatures. 1 NR NR–clay NR–octadecylamine NR–organoclay

0.5

ln K

0

–0.5

–1 –1.5 2.15

2.2

2.25 1000/T

2.3

2.35

(K–1)

FIGURE 9.7 ln K versus 1/T and calculation of Ea. Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

VULCANIZATION KINETICS OF RUBBER–ORGANOCLAY NANOCOMPOSITES

295

3.2

Heat flow (mW)

2

2.8

2.6

2.4

2 ºC/min 5 ºC/min 10 ºC/min 15 ºC/min 25 ºC/min 50 ºC/min

2.2 200

150

250

Temperature (ºC)

FIGURE 9.8 Heat flow versus temperature for NR at different heating rates. Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

Interestingly, the Ea of the elastomer slightly decreased in the presence of the clay, this effect being more evident in the organoclay nanocomposites. According to these results, the authors concluded that the processing conditions of NR improved in the presence of the organoclay because a lower amount of energy was required to transform the plastic undeformed material to the elastic finished material. Moreover, they observed that the octadecylamine-modified clay acted as excellent filler for NR compounds, improving its cure characteristics. The ability of the model to describe the vulcanization kinetics of the NR and its nanocomposites is shown in Figure 9.6, where a good agreement between the experimental and theoretical curves can be easily observed. In addition, the authors analyzed the effect of the unmodified clay and organically modified silicates on the NR vulcanization reaction by DSC, under both dynamic and isothermal conditions. Figure 9.8 shows the dynamic DSC curves for neat NR, where the heat flow absorbed during the vulcanization reaction is represented as a function of temperature at different heating rates. It was observed that the curing temperature increased as the heating rate was increased. The temperature of the cure peaks (TP) for all of the studied materials at different heating rates were reported in Table 9.4. The unmodified clay hardly varies the TP, however a sensible increase in the presence of organoclay was observed. These results are in concordance with those previously observed from the curemeter studies. The Ea of the curing process was easily estimated by means of both the Ozawa and Kissinger equations: Ea ¼  2:3R

d log q dð1=TP Þ

ð9:11Þ

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VULCANIZATION CHARACTERISTICS AND CURING KINETIC

TABLE 9.4 TP and Ea in the Dynamic DSC Tests Ea (kJ/mol) Material





q ( C/min)

TP ( C)

NR

2 5 10 15 25 50

166.13 182.68 192.12 201.30 207.40 224.11

NR–unmodified clay

2 5 10 15 25 50

163.80 185.39 190.60 200.40 208.22 225.72

NR–octadecylamine

2 5 10 15 25 50

139.09 155.15 169.46 178.17 190.69 209.51

NR–organoclay

2 5 10 15 25 50

127.51 143.38 158.05 171.20 180.50 194.30

Ozawa

Kissinger

102.9

95.2

97.5

89.7

75.5

68.1

72.4

65.2

Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

Ea ¼  2:3R

d logðq=TP2 Þ dð1=TP Þ

ð9:12Þ

where TP is the temperature of the exothermic peak and q is the heating rate. Plots of the temperature peak and its derivative versus heating rate (Ozawa and Kissinger equations, respectively) gave rise to a straight lines, and from their slopes, the Ea of the process was calculated (Figures 9.9 and 9.10). The results are expressed in Table 9.4. The results are in concordance with those obtained from curemeter testing. No significant differences by adding of unmodified clay in relation to pure NR are observed. However, the authors observed a sensible decrease in the Ea when the organoclay was added to the NR. These results confirmed that the nanostructured

CONCLUSIONS

297

1.8 NR NR–clay NR–octadecylamine NR–organoclay

1.6 1.4

log q

1.2 1 0.8 0.6 0.4 0.2 2

2.1

2.3

2.2

2.4

2.5

1000/T (K–1)

FIGURE 9.9 Calculation of Ea from the Ozawa equation. Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

–3.6 NR NR–clay NR–octadecylamine NR–organoclay

–3.8

log (q /Tp 2)

–4 –4.2 –4.4 –4.6 –4.8 –5 2

2.1

2.3

2.2

2.4

2.5

–1)

1000/T (K

FIGURE 9.10 Calculation of Ea from the Kissinger equation. Reprinted with permission from Ref. [37]. Copyright 2003 John Wiley and Sons.

layered clay favored the processing conditions for NR because a lower energetic requirement for vulcanization was needed.

9.6 CONCLUSIONS The final properties of a rubber are strongly dependent on the curing degree. To date, there is a lack of knowledge of the reaction mechanism that takes place in the

298

VULCANIZATION CHARACTERISTICS AND CURING KINETIC

vulcanization process, in particular for sulfur-based systems. This mechanism is still more complicated in the presence of fillers. Thus, contradictory results regarding the effect of organoclay on the rubber vulcanization reaction can be found in the literature. However, the majority of these studies show that regardless of the rubber nature and the vulcanization system used, the organoclay acts as a nucleating agent accelerating the vulcanization reaction. The explanation to this fact is the formation of a zinc coordination complex during curing in which the amine groups of the modifier intercalated in the clay gallery and the sulfur participate. In addition, the torque value sensibly increases by adding the organoclay, due to the formation of new filler-rubber interactions. This indicates that the organoclay has a direct influence on the rubber cross-linking density. A more detailed study of the influence of organoclay is necessary to understand the significant improvement in physical or mechanical properties observed in these systems.

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CHAPTER 10

MECHANICAL AND FRACTURE MECHANICS PROPERTIES OF RUBBER COMPOSITIONS WITH REINFORCING COMPONENTS KATRIN REINCKE WOLFGANG GRELLMANN

10.1 INTRODUCTION A large proportion of the scientific literature deals with a wide variety of rubber–clay composites based on a number of different polymers. The following list is far from complete, but provides a first impression of the vast field of composites: . . . . . . . . . . . .

Natural rubber (NR) [1–13] (see also Chapter 12) Acrylonitrile rubber (NBR) [14–19] (see also Chapter 13) Hydrogenated acrylonitrile rubber (HNBR) [20] Carboxylated acrylonitrile rubber (XNBR) [21–23] Styrene–butadiene rubber (SBR) [14,22,24–29] (see also Chapter 12) NR/SBR blend [30] Isoprene rubber (IR) [31–33] Isobutylene–isoprene rubber (IIR) [34,35] Chloroprene rubber (CR) [6,36] Polydimethylsiloxane (PDMS) [37,38] Ethylene vinyl acetate (EVA) [39] EPDM [40–42] (see also Chapter 15)

As described in detail in previous chapters, naturally appearing clays cannot be dispersed very well in a polymeric matrix because of their hydrophobic character. Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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Furthermore, because of the small distance between the clay layers in unmodified clay, the intercalation/exfoliation process is prevented. In general, it is assumed that clays must be treated via ion exchange to get compatible fillers for polymers. This process is called organophilization, and many efforts have been made to successfully trigger it in clay. However, in Ref. [43], a study was reported that showed the effect of a swelling agent on the increase in the interlayer space (d spacing or formerly also known as “intergallery” distance/space) of montmorillonite (MMT) without organophilization. The dispersion increased and a certain amount of intercalation could be obtained. Both the incompatibility between polymer and clay and the lack of intercalation lead to a morphology with filler agglomerates, which is often associated with a lower level of mechanical properties [4,34–36]. Therefore, it is usually assumed that good filler dispersion is one of the basic conditions for sufficient reinforcement effects and is strongly connected to the surface characteristics of the filler. This means that the type of organophilization also plays an important role in reinforcement. Furthermore, due to the intercalation of polymer chains, hydrodynamic reinforcement is increased [31]. Beside the dispersion of the filler particles, for polymer/clay compounds the state of structural degradation of the filler must also be optimal. The first step to obtain an intercalated filler structure is the increase in the distance between the clay layers (interlayer distance); the second step is the exfoliation of the layers, which means the distance is so large that an ordered structure no longer exists. Although it is not always easy to prove the structure status of clay fillers reliably, many papers, for example, [8,9,17,20,28,33,34], discuss a correlation between the degree of intercalation and/or exfoliation of the filler and the mechanical properties. In these papers, intercalated, intercalated/exfoliated, and fully exfoliated filler structures were described. The production of rubber/clay materials is, similar to traditional rubber materials, a very complex process. Many factors such as the number of revolutions during mixing, mixing temperature and mixing time, and time of adding the ingredients of the mixture can influence the morphology and the physical properties of the vulcanizate [21,40,44]. However, the influence of the processing conditions on the mechanical properties is not the focus of this chapter. In general, the aim of material development and optimization is to enhance, or at least to maintain the material properties such as mechanical properties like strength, stiffness, hardness, and toughness, and flame retardance, gas barrier properties, and thermal properties. The mechanical properties especially encompass a wide variety of parameters, which can be determined to characterize different material properties under various loading conditions. This should be taken into account for the selection of certain material testing methods because due to the viscoelasticity of elastomers, time and temperature effects play an important role. However, in many publications, the only test method used is the tensile test. This chapter therefore aims to show the importance and validity of other methods of mechanical testing. For this purpose, important methods of elastomer testing and diagnostics were selected. Each of the sections begins by providing a brief description of the method and experiment,

TESTING OF VISCOELASTIC AND MECHANICAL PROPERTIES

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followed by discussing test results, if available. However, the examples do not always deal with elastomer/nanoclay compounds, but alternatively results from traditional materials are presented. In addition, the sections contain information got from the literature. More experimental details for each test method can be found, for example, in Ref. [45] or in the related standards.

10.2 TESTING OF VISCOELASTIC AND MECHANICAL PROPERTIES OF REINFORCED ELASTOMERIC MATERIALS One important aspect of mechanical testing is specimen preparation. However, because it is similar for all mechanical tests, only a short discussion is required at this point. Specimens for elastomer testing can be punched from vulcanized plates by using cutters having the necessary form and dimension. It is important that these cutters do not have irregular edges because of a possible notching effect. For example, in the case of tensile tests, such irregularities could lead to prior failure of the specimen. Furthermore, a certain sharpness of the cutter must be ensured. If a defect at the elastomer plate is visible, the specimen should be taken from another area. In addition, a certain distance of the specimen punching to the borders of the plates should be kept. Ideally, the specimens should be taken from the same direction. When notches are to be introduced, a sharp metal blade should be used. Recent systematic investigations [46] showed that there is a material-dependent influence of notching on crack propagation. From these results, it was concluded that the variation in the notching angles leads to different stress states in front of the notch and therefore the test results vary. Also, a reproducible notch depth must be realized. 10.2.1 Dynamic–Mechanical Analysis The dynamic–mechanical (thermal) analysis (DM(T)A) is one of the most important instruments that can be used to quantify the viscoelastic behavior of polymers. In general, the cyclic, sinusoidal loading of a specimen can be realized as a tension, pressure, or torsional loading. Such tests can be performed as a temperature or a frequency sweep. Due to the viscoelastic nature of the material, a phase shift between stress s and strain e appears with a certain angle d. Therefore, to describe the connection between s and e, the complex modulus E
is introduced, which has a real part E0 and an imaginary part E00 : E * ¼ E0 þ i E

00

ð10:1Þ

with E0 denoting the storage modulus and E00 the loss modulus. From the relation of E00 and E0 , the loss factor tan d can be determined. More details can be found in Refs. [45,47]. In general, in dependence on temperature, four typical areas can be distinguished for the storage modulus (see Figure 10.1). Regarding the application of a material, the temperature of the glass transition area plays a particular role because at this polymer-specific value, the material changes its mechanical properties

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(a) IV

III

II

Storage modulus log G ′; log E ′

I

Tg Temperature (T ) 1.0 (b) 0.8

SBR

tan δ (–)

NR 0.6

EPDM

0.4 0.2

0.0

–100

–50

0

50

100

150

Temperature (ºC)

FIGURE 10.1 Schematic representation of the behavior of the storage modulus (G0 or E0 ) of an elastomeric material as a function of temperature with I – energy-elastic area, II – glass transition area, III – entropy-elastic area, and IV – area of decomposition [47] (a) and tan d for three polymers with each 40 phr carbon black (b).

significantly. The mechanical loss factor tan d shows a maximum value, which corresponds to the glass transition temperature Tg and depends strongly on the type of the polymer. By adding fillers, the level of the storage modulus in the entropy-elastic range is usually enhanced due to reinforcement effects. Another important point is the height of the tan d peak, which should also be connected to the reinforcement level. The higher the tan d in the main relaxation area, the lower the reinforcement efficiency [36]. The increase in storage modulus, at least within the entropy-elastic range by the addition of clays in comparison to the neat rubbers, was often reported [7,11,15,17,22,24,36]. Furthermore, the increase in filler loading usually led to an increase in the storage moduli of the materials [4,17,20,21,24,25,32].

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Liu et al. [42] showed that the storage modulus above glass transition can be enhanced by increasing the content of maleic anhydride (MA) used as compatibilizer. The MA acts both as a promoter of intercalation and as a cross-linking agent. Partly, the glass transition temperature is influenced by addition of clay to a rubber matrix. In Ref. [28] an increase in the glass transition temperature Tg of SBR/clay compounds was described when modified clay filler was used. As a reason for this behavior, restricted polymer chain mobility within the filler layers was given, which is also a sign of the intercalation of polymers. For NR/clay, NBR/clay, HNBR/clay, and IIR/clay vulcanizates, also an increase in Tg by adding organically modified montmorillonite (OMMT) was reported [3,15,20,33]. Here, the shift in Tg depended on the use of silane as compatibilizer, the production method, the use of stearic acid, the clay content, and the type of the clay modifier. Because the increase in Tg is caused by a kinetic trapping of the polymer chains, as it was mentioned in Ref. [48], conclusions can be drawn regarding the degree of polymer–filler interaction when such an effect is observed. For S-SBR vulcanizates with OMMT and partly additional silica as fillers, a varying Tg was observed when changing the filler content too [25]. However, the authors stated that this shift of max. 4 K cannot be attributed to the filler. Besides the shift of the maximum in tan d, the form and the height of the tan d peak can be influenced by the addition of clay. If interactions at the filler interface take place, leading to an immobilization of polymer chains, the peak becomes narrower and its height is reduced. [3]. Varghese et al. [11] reported the appearance of a shoulder within the peak. It was suggested that the reason for this shoulder is the development of different parts of polymer molecules with varying chain mobility. However, in some cases, no influence of the filler addition on Tg was observed [36,40]. When DMA experiments are performed at a constant temperature and increasing amplitude of oscillation (amplitude sweep), the filler–filler interaction can be characterized. In principle, different components contribute to the height of the modulus. A certain contribution to the modulus arises from the polymer network and a further one, called hydrodynamic effect, depends strongly on the volume fraction of the filler and its shape factor. The latter is the ratio of the longest to the shortest dimension of the filler. For the increase in the modulus of an elastomer, the contribution stemming from the polymer–filler and filler–filler interactions (filler networking) are very important. The storage modulus of a filler-reinforced elastomer material as a function of the deformation amplitude can be characterized by a decrease at higher frequencies. According to the theory, with increasing amplitude, the filler network breaks down and, therefore, a decreasing modulus can be observed. This is known as the Payne effect [49]. Das et al. [36] reported DMA experiments with an amplitude sweep for CR rubber with different nanofillers. It was found that the level of the real part of the shear modulus G0 and therefore the reinforcing efficiency is different, but independent of the kind of filler, which means no Payne effect took place. This was ascribed to the low filler content of 5 phr. Similar behavior was reported both for NBR rubber vulcanizates reinforced with 5 phr organoclay and for comparable materials with a varying cross-linking system [15]. In contrast, in Ref. [7] for NR–organoclay, at each used organoclay content, a decrease in G0 with

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rising shear amplitude was reported, indicating a stronger promotion of nonlinear behavior due to the organoclay compared to traditional fillers. 10.2.2 Tensile Testing The conventional tensile test is one of the most important methods of polymer testing. Also for elastomers, it is one of the basic tests for determining the quantitative characterization of strength and deformability, which are often described as “ultimate properties.” The instructions for the performance of tensile tests for elastomers are summarized in different standards, for example, the German standard DIN 53504 [50], the ASTM D 412 [51], and the international standard ISO 37 [52]. Usually, dumbbell specimens should be used, which are loaded in tension by using a universal testing machine and test speeds between 50 and 500 mm/min. During loading, the force is measured with a suitable load cell and the elongation of the specimen is measured preferably with a displacement transducer in the parallel part of the specimen. On the basis of these measuring values, the stress–strain diagram (s–e diagram) is calculated and can be analyzed. The strength of the material is normally expressed by the tensile stress at maximum (tensile strength) smax. In most cases, the maximum is also the point of fracture. Therefore, the belonging deformation is called tensile strain at break eR and is a measure of the material’s deformability. For the material behavior between the initial loading and the fracture, stress values at different strains can be determined, for example, s100, s200, or s300. These values are often indicated as “moduli”; however, this is not the proper denotation. Furthermore, the determination of the work to fracture Wb that corresponds to the area under the stress–strain curve can complete the results of conventional tensile tests, but it is important to note that Wb is not a measure of material toughness. It may correlate with a toughness value, but not in every case. The form and the characteristic parameters of a s–e diagram strongly depend on the material investigated; therefore, structural changes, for example, as a consequence of variation in the rubber mixture, can be quantitatively registered. As an example, natural rubber should be mentioned, which shows the well-known straininduced crystallization. In contrast to a noncrystallizing material, a typical s-shape of the s–e diagram follows for NR with a large increase in the slope of the diagram at higher deformation due to the self-reinforcing effect of the strain crystallization. Furthermore, with the addition of a filler or with the variation in the cross-link density of the material, the tensile strength and the tensile strain at break can be influenced. Figure 10.2 provides examples of stress–strain diagrams, tensile strength smax, and tensile strain at break eR of IR/organoclay vulcanizates for IR/OMMT materials with different filler contents. It can be seen that at lower deformations of up to approximately 400%, the filler content exerts no influence. The effect of the filler on the stresses at 100, 200, and 300% strain are comparable. This means that from these values no reinforcement effect of the rising filler content can be derived. However, the strength and strain increase with increasing content of Dellite 67G. The conclusion from the literature is that in almost all cases the tensile stress at a given strain, for example, s100 or s300, increases due to the addition of modified or

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40

311

(a)

σ (MPa)

30

20

22 phr LS 7 phr LS 2 phr LS

10

0

0

200

400

600

800

ε (%) 35

800

(b)

700 25

εR (%)

σmax (MPa)

30

20 600

σmax εR

15 10 0

5

10

15

20

500 25

Content of Dellite 67G (phr)

FIGURE 10.2 Results of tensile tests according to DIN 53504 for IR compounds with Dellite 67G. (a) Examples of engineering stress–strain diagrams (s–e diagrams); gray circles mark the point of fracture, where tensile strength smax and tensile strain at break eR are determined; lines mark different levels of deformation at which usually tensile stresses are determined, here s100, s200, and s300. (b) Tensile strength smax and tensile strain at break eR of the IR compound depending on the content of Dellite 67G.

unmodified clay filler [4,7,8,11,17,21,22,24,36]. How pronounced this increase is in comparison to the neat material depends on certain conditions, which are discussed more or less in detail in the individual papers. In general, different reasons for the better mechanical performance expressed by s100, s300, or smax are given. These are, for example, a better dispersion of the filler due to increased compatibility [17,24,34], more pronounced polymer–filler interaction [36], and the state of intercalation/ exfoliation due to modification of the clay [8,11,36]. According to Ref. [25], the decrease in strength and deformability observed for the SBR materials with higher filler content can be explained by the filler agglomeration, which means a worse dispersion.

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In a study by Ramorino [1], the initial modulus of NR/OMMT materials was determined and discussed concerning the OMMT content, the test temperature, and the test speed. It was shown that the stiffness of the vulcanizates increases with the filler addition and the test speed and decreases with temperature, indicating the viscoelastic character of the underlying processes, especially within the area of the polymer–filler interface. Furthermore, a percolation threshold at a filler content of 8 phr was observed. This means that filler networking is apparent, leading to higher stiffness of the material. The material parameter tensile strength may react to the clay addition with an increase [6–8,11,17,21,24,34,35] or a decrease as it was reported in Refs [4,13] for NR compounds with pristine MMT or OMMT, respectively. According to Ref. [4], this was to be attributed to poor adhesion of the filler, while in Ref. [13] the reason was assumed to be the disturbance of the strain crystallization of the NR. In contrast to the strength properties, the deformability of rubber/clay compounds is often reduced in comparison to the neat material or to materials with traditional fillers [4,11,13,17,21,25,34]. However, there are also examples where the strain at break maintains or increases with clay addition. For example, this behavior was reported for SBR/clay [24], NR/clay [7–9], and XNBR/clay [21] vulcanizates. Often maxima of the mechanical parameters are obtained with increasing filler content as it was described for s100 and smax of an IIR/OMMT material [34] where s100 and the tensile strength smax were reduced with filler amounts larger than 10 and 5 phr, respectively. Also for XNBR, NR, and SBR, an optimum in tensile properties including tensile strength and/or tensile strain at break was found to depend on the clay content [4,6,12,22,25,26]. As a reason for the decreasing tensile properties, the authors [26] stated that the stress concentrations appearing under mechanical loading at the edges of the layer stacks lead to failure, which is more pronounced at higher filler concentration. Furthermore, with increasing amounts of clay, it becomes more difficult to exfoliate it and filler aggregation is assumed, leading to a reduction in tensile strength. Also, the authors of Ref. [4] explain the decreasing tensile strength at higher clay loading with filler aggregation. Another aspect described in the literature is that when good interfacial adhesion between filler and polymer exists, the layer bundles may separate during loading. When it is assumed that the new surface does not have a connection to the polymer, these areas act as defects and the property level may be reduced. Some tensile test results are reported to show a more or less pronounced increase in tensile strength of rubber materials with unmodified clay. One example is given in Ref. [35] for IIR/clay compounds. The unmodified clay gave an increase in tensile strength up to approximately 13 MPa with the highest clay loading of 20 phr. In comparison, the material with the organoclay reached a tensile strength of approximately 15 MPa with 20 phr filler. For a CR material with unmodified MMT, also an increase in smax compared to the neat CR was reported [36]. More interesting results were published by Hrachova et al. [9], who investigated NR materials reinforced with natural clay as well as with variously modified clays. It was found that addition of the unmodified clay (up to 10 phr) causes an increase in tensile strength and in tensile strain at break. At the same time, the stress at 100% strain

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showed a general decrease. Furthermore, the authors investigated vulcanizates for which additional 15 phr of silica were used. It was illustrated that the most pronounced increase in tensile strength and tensile strain at break at medium clay content was obtained for the compound with a combination of unmodified clay and silica. Similar results were presented by the same authors in Ref. [10], where NR compounds with a combination of different clays and silica were investigated. Both the strength and the deformability are strongly increased. This is the case not only for organomodified clays but also for neat clay. On the basis of X-ray diffraction (XRD) measurements, the interlayer distance was calculated for these materials. The modification of the clays led to an increased d spacing, but only in materials without silica. When silica was added, all four MMTs had a similar interlayer distance in the range of 1.43 and 1.45, respectively. This means that the effect of the modifications seems to have disappeared. Similar observations were made by Varghese et al. [11], who reported not only an intercalation/exfoliation but also a confined structure from decreased interlayer distance caused by different polymer chain mobility in various areas. 10.2.3 Assessment of Toughness Behavior under Impact-Like Loading Conditions The toughness behavior of an elastomeric material can be investigated by using the conventional tensile impact test, which is standardized [53]. The aim of the test is to characterize the behavior of specimens under relatively high impact velocity and to assess the toughness or brittleness of the material, respectively. Toughness is defined here as the ability of the material to consume energy before the fracture appears. For performing conventional tensile impact tests according to method A in Ref. [53], commercial impact pendulum devices with working capacities of up to 50 J and with specially formed hammers, a clamping device, and a specimen support are used. According to method A in Ref. [53], specimens with V-notches on each side (type 1) or dumbbell specimens without a notch (type 3) can be used. The notches of type 1 specimens are each 2 mm in size and have a relatively large notch tip radius of 1 mm. Depending on the type of the pendulum device and the pendulum hammer used, the test speed is in the range of approximately 2.9 m/s or approximately 3.7 m/s. During the test, the impact energy is determined. Afterward, the tensile-impact toughness atU or the notched tensile-impact toughness atN is determined using the following equation: atU

or atN ¼

Ec BðW aÞ

ð10:2Þ

with Ec denoting corrected impact energy, B specimen thickness, W specimen width, and a initial notch depth. One example of such results is given in Ref. [54], where the notched tensileimpact toughness of filler-reinforced SBR materials was presented as a function of the filler content. As it was shown, the toughness increases with filler loading, but the

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standard deviation is relatively large. Therefore, it is difficult to derive reliable statements regarding the influence of the filler type on the energy consumption capability, especially at higher filler contents. Another experimental possibility to characterize the material behavior at higher loading speeds is the instrumented free falling dart test by which the multiaxial deformation behavior can be quantified. Because it is a relatively new field for elastomers, no official test standard exists for this group of materials and the pool of experience is still limited. However, a current standard for thermoplastic testing [55,56] may be adopted for the test procedure. According to [55], round or rectangular, flat specimens with a minimum diameter or edge length of 65 mm, respectively, should be used. For the test, the specimen is clamped between the support and an upper clamping ring and loaded by the falling dart perpendicularly to its surface (see Figure 10.3). In this way, a multiaxial stress state within the specimen emerges when the falling dart deforms the specimen. During the test, a load–time (F–t) diagram is recorded from which the load–deformation (F–l) diagram is calculated, which can be analyzed regarding the determination of maximum load FM and maximum deformation lM, as well as the corresponding energy EM, for example (Figure 10.3). When a temperature chamber is available, the deformation behavior can be investigated also at varying test temperature; this should be of practical interest for viscoelastic materials. Furthermore, most falling dart testers allow for changing of the test speed, so that experiments concerning the influence of the loading speed can also be performed. One example of the application of the instrumented free falling dart test was presented in Ref. [57]. Here, results ofinstrumented free falling dart tests for NR/carbon black (CB) compounds with different CB contents were described. The tests were performed in dependence on the test temperature. It was found that the load level increases with the CB content, as expected. At  25 C a maximum of the maximum load FM seemed to appear, which may correlate with an increase in Tg due to the higher loading speed (time–temperature superposition principle), or may reflect a temperature-dependent effect of the strain crystallization of the NR. The deformability expressed as lM was little different. In general, the deformability changes much more with temperature than the maximum load does. An increasing test temperature

Spherical calotte Clamping ring Specimen Specimen support Basis Damper

FM Load

Piston Load measuring (strain gauge)

FP ER

Eu lM

lP

Deformation

FIGURE 10.3 Test principle of an instrumented free falling dart test [55] and schematic load–deformation diagram.

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315

resulted in a rising lM value. In addition to analyzing the test results according to DIN EN ISO 6603-2, slope S of the F–f diagrams in the initial part and the deformation at a load of 300 N d300 were determined for the NR compound with 52 phr CB to obtain more information on the material’s behavior. Different from FM, slope S showed no maximum, but a continuous decrease, similar to storage modulus E0 . 10.2.4 Hardness Testing For elastomers, hardness tests are basic tests since a certain degree of hardness is necessary for the specified application of the material in most cases. Very often, Shore A hardness is determined. Another well-known method for testing elastomers’ hardness is the IRHD testing. In general, with hardness testing, the material’s resistance against the indentation of a harder body is quantified. There are a number of standards for the performance of Shore A or D tests [58,59], as well as IRHD tests for elastomers [60,61]. Usually, Shore A testing is used for materials with a mean hardness, while shore D testing can also be used for “harder” materials. Shore tests should be performed with 6 mm thick specimens, so that an influence of the background can be excluded. In cases where the sample is not large enough, which is often the case, when, for example, testing components such as seals, microhardness testers can be used. Because hardness testing is easy to perform, other important material parameters like strength or abrasion resistance are often predicted on the basis of hardness values. However, such correlations do not exist generally and for product design and material development, hardness values cannot be used alone. In Figure 10.4, the tensile strength values smax and the related Shore A hardness values of a number of fillerreinforced elastomers are plotted. As it can be seen, no correlation exists between both mechanical parameters. 40

Tensile strength σmax (MPa)

Elastomers with layered silicate Elastomers with CB or silicate 30

20

10

0 30

40

50

60

70

80

90

Shore A (–)

FIGURE 10.4 Shore A hardness and tensile strength of different elastomeric materials (EPDM, SBR, NR with carbon black or silica, as well as NR with different layered silicates).

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MECHANICAL AND FRACTURE MECHANICS PROPERTIES

TABLE 10.1 Shore A Hardness Values of IR/Dellite 67G and NR/Nanofil Compounds Filler Content (phr)

Shore A at 23 C

Shore A at 70 C

2 7 22 0 5 10 15 60 70

42.7  0.5 44.3  0.3 51.7  0.4 29.4  0.1 31.7  0.0 34.1  0.8 37.0  0.4 57.2  0.1 60.3  0.3

29.9 31.9 34.3 36.5 53.5 56.3

IR/Dellite 67G

NR/Nanofil

In contrast, a correlation between the material composition with polymer, filler, plasticizer, cross-linking system, and so on, and the resulting hardness of the material is given; however, the sensitiveness of the hardness values regarding structural changes must be taken into account. For example, during a long-term thermal–medial loading of two EPDM elastomers, the Shore A hardness changed by 1 and 3%, respectively, while other material properties such as tensile strength or tear strength were reduced by up to 46% [62]. Furthermore, in an investigation by Rattanasom and Prasertsri [63] of NR materials with a filler combination of calcinated clay and carbon black, all materials showed the same hardness level, but other mechanical parameters such as tensile strength or tensile strain at break differed. Table 10.1 shows results of Shore A hardness measurements for different rubberlayered silicate materials. Both types of layered silicate are organically modified and for the NR/Nanofil compounds, an intercalated/exfoliated structure was observed by TEM images [13]. From Table 10.1, it can be seen that increasing filler content always leads to an increase in the material’s hardness, but the type of filler or the combination polymer–filler, respectively, is important in determining the level of hardness. In Ref. [35], an increase in Shore A hardness values of IIR/clay vulcanizates produced by melt or solution intercalation with the content of clay filler was also reported. The material from the solution intercalation showed a higher level of mechanical properties like tensile strength and hardness compared to the melt intercalated material. Further publications showed an increase in the Shore A hardness of rubber/clay vulcanizates with different polymer matrices such as NR [6,11], IIR [36], and CR [6] due to the addition of modified or unmodified clay and/or an increasing filler content. 10.2.5 Special Methods 10.2.5.1 Tear Behavior In the elastomer industry and technology, one widespread testing method is the tear test, which can be performed with a variety of specimens and test conditions. By means of this test, the behavior of a material in the

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317

presence of a sharp cut is characterized. There are different standards, for example [64,65], which can be applied to determine the tear resistance, also called tear strength, of an elastomeric material. Often, crescent or trouser specimens are used. Such tests can be done by using universal testing machines. The loading speeds can range between 50 and 500 mm/min. During a test, the load–displacement diagram is recorded, from which maximum load Fmax or the median of load Fmed can be obtained. According to Ts ¼

F B

ð10:3Þ

where B denotes specimen thickness, tear strength (Ts) can be calculated. Ts is strongly influenced by the material composition, but it is possible to make only a few general statements. For example, the type of the polymer used determines the kind of crack propagation. Strain crystallizing materials may show knotty tearing or stick slip behavior, which results mostly in higher loads until the specimen is completely torn. On the other hand, especially EPDM is known for its notch sensitivity. Therefore, the tear strength of such materials is expected to be lower than for NR, for example. Also, the filler type and content or the cross-link density can have a great influence on the tear strength. Examples are provided in Figure 10.5, where tear strength Ts, determined according to Ref. [64] by using trouser specimens, is shown in dependence on the clay content for two different material series: IR/ organoclay and NR/organoclay. Increasing the organoclay content results in a maximum value of Ts at a filler content of 7 phr. Adding Nanofil to the NR, the tear resistance decreases with lower filler contents and reaches the initial level with 60 and 70 phr, respectively. The two series in Figure 10.5b should illustrate that no significant influence of the direction of the specimen preparation exists because the specimens for the second series were taken from the plate rotated by 90 . Interestingly, the comparison of the level of the tear strength values of the IR/and the NR/ organoclay materials showed no such strong differences as it is the case for the tensile strength, tensile strain at break (Figure 10.2), and Shore A hardness (Table 10.1). Besides the results of this study, a number of papers by other authors also discuss not only the tensile properties of the materials but also their tear resistance. For example, an increase in tear strength Ts with (increasing) clay addition was described for NR in Refs [6,11], for NBR in Refs [17,19], and for SBR in Ref. [24]. Often this increase was attributed to a better dispersion of the filler. Furthermore, in Ref. [19] the influence of the organomodification of sodium-MMT (Na-MMT) on the tear behavior was discussed. The authors found a significant difference in tear strength between the unmodified and the modified type of the clay. SEM investigation proved that the filler of the NBR/MMT compound is not really connected to the polymer matrix, leading to the lowest mechanical property level including the lowest tear strength value. 10.2.5.2 Determination of the Compression Set For gaskets, for example, it is important to know how materials retain their elastic properties after a long-term

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10

(a)

Ts (N/mm)

8 6 4 2 0

25

20

15

10

5

0

Content of Dellite 67G (phr) 10

(b)

Ts (N/mm)

8 6 4 Along Cross-wise

2 0

0

10

20

30

40

50

60

70

80

Content of nanofil (phr)

FIGURE 10.5 Tear strength Ts of IR/clay (a) and of NR/clay (b) compounds from tear tests.

mechanical and/or thermal–medial loading. To characterize this application-oriented material property, usually the compression set (or permanent set) is determined. Compression set is a measure of the ability of the material to make an elastic recovery after a more or less long-lasting deformation. The higher the determined value, the lower the elastic re-deformation. Such experiments are realized by applying a defined deformation on a round specimen with a certain diameter and height. Other test conditions such as time of loading, level of deformation, temperature, and so on can be fixed according to ISO 817 [66]. The test is completed by releasing the compression load, and 30 min later, the remaining height of the specimen is measured. From the relation between initial and after-test height, the compression set is calculated. Rubber–clay composites are an interesting material group also for their application as gaskets, for example, because of their good permeation behavior. For this reason, some recent publications reported the compression or permanent set of such

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materials. For example, in Ref. [36] it was stated that the compression set of a CR rubber with 5 phr MMT is at the same level as with 5 phr N220, while that of the OMMT material is increased. However, it should be mentioned that other material properties such as storage modulus G0 , tensile strength smax, and Shore A hardness are on a higher level for the OMMT material. In contrast, in Refs [5,8] a lower level of the permanent set was obtained with OMMT compared to the MMT addition. The authors in Ref. [39] showed that the use of silane resulted in a lower compression set compared to the materials without silane. It was concluded that the stronger chemical bonding of the polymer chains to the filler in the case of silane usage reduced the chain slippage at the filler surface, which is considered as one main reason for the loss of elastic recovery ability, expressed, for example, as the compression set. Therefore, it was stated that an increased polymer–filler coupling should result in a better compression set. However, other results from an NR material with 40 phr of an active carbon black showed a comparable high compression set of this material [8], although here usually a sufficient polymer–filler interaction was assumed. This means that other mechanisms must contribute to the elastic recovery behavior of an elastomer material. 10.2.5.3 Fatigue Behavior Another important material property of elastomers is their fatigue behavior. However, only very few studies dealing with such test results for rubber/clay materials have been published. For example, in Refs [29,30] the flex fatigue test of SBR vulcanizates with a filler combination of CB and clay and of NR/ SBR/clay compounds was used. This method, also called DeMattia test, is a relatively simple test and can be used to characterize the fatigue behavior. In general, such tests can be performed according to Ref. [67] with relative thick strip-shaped specimen with a groove in the middle. The specimens are cyclically completely folded, so that the maximum deformation appears at the tip of the groove. The tests are usually performed with a frequency of 5 Hz at room temperature. The aim of the test is to determine the number of cycles possible before crack growth or the number of cycles necessary until a crack of a certain size has grown. The results in Ref. [29] showed that due to the substitution of 2–5 phr of CB by Na-MMT, the flex fatigue life was improved, especially with 4 and 5 phr clay. The authors stated that this should be attributed to the high aspect ratio of the nanodispersed clay layers. Cataldo [30] reported a relatively non-influenced flex fatigue life of the materials when the clay content is increased by up to 15 phr. This result correlates in this special case with the tear resistance, which also showed no significant changes with increasing clay content.

10.3 CHARACTERIZATION OF THE FRACTURE BEHAVIOR OF ELASTOMERS 10.3.1 Fracture Mechanics Concepts During the past 90 years, different fracture mechanics concepts have been developed with the aim of a failure-related description of a material. Based on the well-known

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MECHANICAL AND FRACTURE MECHANICS PROPERTIES

Griffith criterion, at first the concepts of linear-elastic fracture mechanics (LEFM) were applied. LEFM can be applied for the characterization of materials showing linear-elastic deformation. However, the deformation of many materials is characterized by a plastic deformation before the unstable fracture occurs. Therefore, the LEFM was extended to the LEFM with small-scale yielding that considers a small area of plastic deformation in front of a crack tip, called the plastic zone. The application of concepts of LEFM allows determining stress intensity factor K or energy release rate G necessary for the quantitative description of the stress field in front of a crack tip. Here, the stress intensity factor is mostly used in experiments because of its proportionality to the applied stress. In addition, K describes the stress field around the crack tip independent of the Young’s modulus of the material. Therefore, the primary fracture criterion of LEFM is the critical stress intensity factor called fracture toughness Kc, for which an inactive crack begins to grow in an unstable manner. The fracture toughness is a geometry-independent material parameter if a plane-strain stress state is given. Another way of obtaining a fracture criterion of LEFM is to use a modified energy balance. Here, it is assumed that for the propagation of a crack certain energy is necessary, which comes either from the applied load or from the release of elastic deformation energy in a deformed body. As a result, the critical energy release rate Gc is determined as a material parameter describing fracture resistance [68,69]. With proceeding material development and optimization, fracture theories have also been advanced, and concepts of elastic–plastic fracture mechanics (EPFM) such as crack tip opening displacement (CTOD) and the J-integral concept expanded into material science. They have reached their greatest importance for (crack) toughness characterization because they can be applied to materials with extensive plastic deformation in front of the crack tip. The J-integral as fracture criterion has highest importance owing to its energetic interpretation of the fracture process. For an experimental determination of J, different approximation solutions are available [70]. Furthermore, unstable and stable crack growth must be characterized separately. A crack is denoted as “unstable” when its speed is very high and the crack propagation takes place with energy release. In contrast, a “stable” crack is characterized by a comparatively low crack speed; energy is dissipated during this process, which can be stopped by taking away the external loading. Considering a material showing stable crack propagation, the crack resistance curve (R-curve) concept of EPFM is appropriate to describe different states of a stable fracture. An R-curve represents the relation of a loading parameter (e.g., J) and a damage parameter (e.g., stable crack growth Da). For obtaining such data, crack resistance curves can be recorded by using single-specimen or multiple-specimen methods under quasistatic or impact-like loading conditions. However, the conditions for the application of the aforementioned fracture mechanics concepts are not met for elastomers in all aspects because they show neither linear-elastic nor elastic–plastic deformation. Furthermore, the deformation during mechanical loading is not limited to the area in front of the crack. This makes it difficult to transfer these concepts directly to this group of materials. On the basis of the Griffith criterion, Rivlin and Thomas made an energetic interpretation of the

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fracture process of elastomers in 1953 and introduced the tearing energy T as a fracture parameter for elastomers [71]. It was assumed that energy dissipation occurs next to the crack tip independent of the specimen geometry used. According to Rivlin and Thomas’s definition, energy release rate G, the J-integral, and tearing energy T are formally identical (see Eq. (10.4)) and therefore these parameters can be considered as equivalent. G¼J¼T ¼

1 @U B @a

ð10:4Þ

with U denoting potential energy, B specimen thickness, and a crack length [45,68,69]. 10.3.2 Experimental Methods 10.3.2.1 Fracture Mechanics Tests under Quasistatic Loading The experimentally easiest way to obtain fracture mechanics parameters is to do quasistatic tests by using universal testing machines. The experience of the last years [13,54,72,73] showed that crack resistance curves determined by the application of a single-specimen method can deliver valuable information about the fracture behavior of elastomers related to the crack initiation and propagation of a stable crack. For such R-curve tests, single-edge-notched tension (SENT) specimens are prepared with certain dimensions and an a/W ratio of at least 0.2. The authors’ own results presented in this chapter were obtained with specimens of width W ¼ 25 mm, length L ¼ 100 mm, and thickness B ¼ 6 mm. The notches must be produced by using a thin metal blade. The specimen thickness should not be too small to exclude or at least to reduce a plane stress state within the specimen. Depending on the material under study, the test speed is fixed in the range of up to 50 mm/min. Before starting the test, the notch tip is prepared with TiO2 powder to enable the examination of the crack initiation point. During the test with monotonically increasing loading of the specimen, the notch tip is observed and pictures of it are taken at different stages of the deformation and tear process. After the crack initiation, the distance between the edges of the notch, defined as crack opening lR, becomes larger with the deformation of the specimen. Investigations indicated that a linear connection between the real crack size Da and lR exists, so that the latter can be used as x-value of the R-curve. The related y-value is obtained on the basis of the recorded load–displacement (F–l) diagram. Deformation energy U is determined by integrating the F–l diagram up to the certain times. The time intervals must correspond to the times of taking the pictures. Afterward, according to J¼

hU BðW aÞ

with h denoting geometry function, the J values are calculated.

ð10:5Þ

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MECHANICAL AND FRACTURE MECHANICS PROPERTIES

When the J–lR data pairs are plotted as a crack resistance (R-) curve, a typical behavior can be observed. From such R-curves, two fracture parameters can be generally obtained. The first one is crack initiation value Ji (sometimes also denoted as JIc), which is determined through observation of the notch tip, and therefore a subjective influence cannot be excluded. The second is a crack propagation value coming from the slope of the R-curve in the first part. Long experience with thermoplastic materials showed that mostly the crack propagation resistance is much more influenced than the crack initiation behavior [45,70], and also for elastomers this phenomenon was proven [74]. In Figure 10.6, R-curves of fillerreinforced EPDM vulcanizates are given as examples. Other results of this study can

J (N/mm)

60

(a)

40 Aromatic content of the oil (%)

2 4

0 3 5

20

0 0

20

10

40

30

60

50

70

Crack opening (mm) 30 (b) 25

5 phr 70 phr

J (N/mm)

20 60 phr 15

10 phr 10 15 phr 5 0 0

10

20

30

40

Crack opening (mm)

FIGURE 10.6 Examples of R-curves of EPDM vulcanizates with white fillers and different mineral oils (a) and of NR/Nanofil compounds with various filler contents (b).

CHARACTERIZATION OF THE FRACTURE BEHAVIOR OF ELASTOMERS

323

be found in Ref. [75]. To investigate the influence of the aromatic content of the mineral oil used as a plasticizer on the mechanical and fracture properties, different grades of such oils were incorporated. The R-curves in Figure 10.6a demonstrate a significant influence of the materials’ structure on the crack propagation behavior. With rising aromatic carbon content of the oil, the cross-linking system was increasingly “blocked”, so that the cross-link density of the materials varied greatly. This resulted in strong differences both in the reached level of J and in the size of crack opening lR before the ultimate tearing occurred. The R-curves in Figure 10.6b also demonstrate a different stable crack propagation behavior. In addition to the influence of the composition, in this case the filler content, on the level of J values, the R-curves of NR/clay materials in Figure 10.6b indicate a transition from stick slip fracture to steady tearing. The latter is characterized by a continuous course of the R-curve, while the stick slip is connected with a pronounced crack opening without a large increase in the level of J, leading to “irregularities” of the R-curve as can be seen for NR with 5–15 phr Nanofil showing some steps with sudden larger increases in the crack opening. As a reason for the changing crack propagation behavior of the higher filled materials, it was assumed that the strain crystallization of the polymer matrix was disturbed by the not fully exfoliated filler particles [13]. The NR material with only 5 phr Nanofil exhibited the best crack resistance behavior. R-curve tests can also be performed in terms of the multiple-specimen method. For this purpose, a number of similar specimens are used to realize for each of them a different loading level after a crack initiation was observed. The loading and unloading cycle is recorded, so that an extended analysis with determination of J and Jdiss (Eq. (10.5)) is possible. Jdiss reflects only the energy determined from the hysteresis loop and therefore it is different from J. After the test, the specimen is completely cut and the size of the stable crack propagation can be measured by using a light microscope. The analysis of the resulting R-curve is similar to that of the single-specimen test as described above. More details can be found in Ref. [76]. 10.3.2.2 Instrumented Notched Tensile-Impact Tests Instrumented notched tensile-impact testing (ITIT) allows for a fracture mechanics toughness characterization of very thin and/or flexible materials on the basis of load–extension (F–l) diagrams. The test setup is similar to that of the tensile-impact test described in Section 10.2.3. However, the pendulum device must be fit with a load cell. According to our own standard [77], double-edge-notched tension specimens (DENT) with the dimensions length L ¼ 64 mm, width W ¼ 10 mm, thickness B ¼ 2 mm, and initial notch depth a ¼ 2 mm are to be used, whose notches where made with a metal blade, so that a small notch tip radius is obtained. Figure 10.7 shows the test configuration with the hammer and a specimen clamped between the crosshead and the fixed clamp. Usually, for each series, the F–l diagrams of 10 specimens are to be recorded and analyzed regarding maximum load Fmax extension at maximum load lmax, energy uptake up to Fmax Amax, and crack propagation energy Ap [78]. As a result, Jd values

324

MECHANICAL AND FRACTURE MECHANICS PROPERTIES

FIGURE 10.7 Instrumented tensile-impact rig of the pendulum device Resil Impactor Junior (CEAST, Pianezza, Italy).

describing the materials’ resistance against crack propagation can be determined according to Eq. (10.6) hAmax ð10:6Þ Jd ¼ BðW aÞ with the geometry function h according to Ref. [69] h ¼ 0:06 þ 5:99


a
a 2
a 3 þ3:29 7:42 W W W

Results of instrumented notched tensile-impact tests, together with results of tear tests for NR–clay composites, are shown in Figure 10.8. It can be seen that the results of the tear tests expressed as the high-speed tearing energy HSTE and tear strength Ts correlate very well, although strongly different loading speeds with 3.7 m/s and 100 mm/min were applied. An increasing Nanofil content does not change the tear resistance. In contrast, the Jd values from the ITIT show a significant increase with the Nanofil content. As the F–l diagrams in Figure 10.8b show, the increase in Jd is due to the increasing load-bearing capacity of the highly filled materials. Although the deformability is reduced by nearly 50%, the maximum load strongly increases and therefore also an increase in Jd occurs. Another finding of the ITIT was that toughness values Jd and measuring values Fmax and lmax of an NR/BR/SBR blend with two kinds of a layered silicate on the basis of bentonite are influenced by the degree of structural degradation of the filler. Similar to the former example, the increase in toughness was not only due to the higher load-bearing capacity but also due to a better deformability, which was derived from the course of Fmax and lmax. Again, mainly the load component of the impact energy increases with the filler content, but not the deformation component. In this case, a load-determined fracture behavior is obvious.

CHARACTERIZATION OF THE FRACTURE BEHAVIOR OF ELASTOMERS

300

(a)

Jd

HSTE

Ts

15 200 10

Jd (N/mm)

HSTE, Ts (N/mm)

20

325

100 5

0 0

20

10

30

40

60

50

70

0 80

Content of layered silicate (phr) 80 (b) 60 phr

Load (N )

60

70 phr

40

15 phr 10 phr 5 phr

20

0

0

20

40

60

80

100

120

Extension (mm)

FIGURE 10.8 Jd values from instrumented notched tensile-impact tests, tear resistance Ts, and the impact tear resistance values HSTE of NR/Nanofil compounds in dependence on the Nanofil content (a), and examples of F–l diagrams from ITIT for each NR/Nanofil compound (b).

When the ITIT should be performed at different temperatures, an external tempering is necessary. For this purpose, special holders for the specimens were developed, so that a fast insertion of a specimen into the pendulum device is possible. Results of such tests are shown in Figure 10.9. It can be seen that the crack toughness behavior of the SBR and the NR compounds differs according to temperature. While the SBR materials showed a maximum at 0 C, the NR did not exhibit such a course. Here, the four materials can be divided into two groups. The two materials with the higher filler content have a constant level of Jd of up to 0 C, then Jd decreases with the temperature. For the unfilled and the material with 20 phr CB, a continuous decrease in Jd with the temperature was detected. Furthermore, the change in the material ranking at different temperatures for the SBR materials was an interesting finding. However, this topic requires further research.

326

MECHANICAL AND FRACTURE MECHANICS PROPERTIES

300

(a) 0 20 phr 40 phr 60 phr

Jd (N/mm)

200

100

0 -80

-40

0

80

40

Temperature (ºC) 300

(b) 0 phr 20 phr 40 phr 60 phr

Jd (N/mm)

200

100

0 -80

-40

0

40

80

Temperature (ºC)

FIGURE 10.9 Influence of the testing temperature on the toughness behavior of SBR/CB (a) and NR/CB compounds (b) for different CB contents.

10.3.2.3 Fracture Mechanics Tests under Cyclic Loading The fatigue behavior of vulcanizates can be investigated with servohydraulic testing machines like the “tear and fatigue analyzer (TFA)” that was developed by Bayer [78] and is commercially available from Coesfeld GmbH, Dortmund, Germany. Many efforts have been made to investigate the quantitative influences of experimental details of such measurements like specimen type or amplitude of the deformation, pre-strain, kind of cyclic loading, and so on. However, experimental problems still remain regarding the production of reproducible test results, as discussed, for example, by Sto
cek et al. [79] and Ziegler et al. [80]. The aim of fracture mechanics tests under cyclic loading is to obtain a crack propagation curve, which is the functional connection between crack velocity da/dN and a fracture mechanics parameter such as stress intensity factor K, the J integral, or

CHARACTERIZATION OF THE FRACTURE BEHAVIOR OF ELASTOMERS

log da/dN

I

II

327

III

da/dN = C (ΔK, ΔJ, ΔT)m

log ΔK, ΔJ, ΔT

FIGURE 10.10 Schematic crack propagation curve according to Refs [81,82].

tearing energy T, which is usually used for elastomers (Figure 10.10). The crack propagation speed corresponds to the change in the crack length Da in dependence on the number of cycles N. The tearing energy is calculated according to the following Eq. (10.7): T ¼ 2kWel aeff

ð10:7Þ

pffiffiffi where k ¼ p= l and l ¼ l=l0 , with l denoting extension, l0 initial gauge length, Wel elastic energy density, and aeff effective crack length. Crack propagation curves can be generally divided into three different ranges. Within ranges I and III, a large influence of the microstructure, the mean stress, and the possible media exists. Most important for the quantitative characterization of the fatigue fracture resistance is range II, in which with double logarithmic plotting a linear connection between the crack propagation speed and the fracture mechanics parameter appears. This can be described by the Paris–Erdogan equation (Eq. (10.8)): da ¼ C  ðDK; DJ; DTÞm dN

ð10:8Þ

with m, C denoting material-specific parameters [81,82]. In general, for rubber/clay materials, hardly any result of fracture mechanics tests under cyclic loading has been reported yet. Figure 10.11 shows results of TFA tests with NR/Nanofil compounds in comparison to NR/silica. These tests were performed with SENT specimens under a pulsed cyclic load of 30 Hz. It can be seen that the crack propagation behavior differs according to the Nanofil content. The slope of the crack propagation curve is decreased by adding 60 phr of this OMMT, which means that the velocity of the crack growth is reduced compared to the material with the same amount of silica and to that with 15 phr Nanofil. However, the scattering of the measuring points is relatively large, so that it is difficult to make a reliable statement.

328

MECHANICAL AND FRACTURE MECHANICS PROPERTIES

da/dN = 0.0012 T

1.74

1E-3 da/dN = 4.51.10 T

da/dN (mm/cycle)

–4

0.25

1E-4 da/dN = 0.0011 T

1.31

60 phr Nanofil 15 phr Nanofil 60 phr Sillca

1E-5

0.01

0.1

1

T (N/mm)

FIGURE 10.11 TFA results of NR/Nanofil silicate compounds and an NR/silica vulcanizate for comparison purposes; test temperature was 60 C.

10.4 MECHANISM OF REINFORCEMENT IN RUBBER–CLAY COMPOSITES It is generally assumed that the improvement in polymer–filler interaction is one of the most important aspects regarding an increased level of the mechanical properties of rubber/filler compounds. Depending on the kind of carbon black, as the most commonly used reinforcing agent for elastomers, its surface area has many active sites on which polymer chains can interact. Here, the polymer–filler interaction is principally given due to the nature of the carbon black. When precipitated silica as the other widely used “conventional” filler is applied, a compatibilizer such as silane must be used. Otherwise, because of the surface characteristics, the filler–filler interaction is too large, leading to the appearance of filler agglomeration and preventing sufficient coupling between polymer and filler. Based on this knowledge, also for elastomeric materials with alternative fillers such as layered silicates, it is assumed that the polymer–filler interaction is the precondition for sufficient reinforcement. In other words, the interface between the filler and the polymer must be optimized. Therefore, many attempts have been made to optimize the polymer and/or the filler surface. The special morphology of clay fillers makes it a little more complicated to discuss the reinforcement mechanisms. There is a great number of influencing factors that are partly counterproductive. Apparently, there is no “rule” for the filler reinforcement with clays. Depending on the polymer used, filler type, compatibilizer, processing, and so on, different results were obtained as it was shown above. The authors’ own results and those described in the literature discussed here indicate that no general correlations between composition and physical properties of a polymer/ clay compound exist; instead, for each polymer/filler system, an optimum must be found. One example of filler surface optimization is given in Ref. [14], where clays

MECHANISM OF REINFORCEMENT IN RUBBER–CLAY COMPOSITES

329

were surface treated with silane, followed by electron-beam irradiation to make them more compatible to SBR and NBR rubber. Although no larger d spacing could be proved due to the surface modification, the mechanical properties, expressed by the results of tensile and tear tests, could be improved, especially for SBR as a non-polar rubber. In general, the behavior of polymer chains in constraint situations as a basis for understanding the often quoted polymer–filler interaction is in the focus of much theoretical research work [48,83,84]. However, this chapter concentrates on more practical aspects of polymer/filler compounds such as compatibilizers, particle size, and so on. In the previous sections, some aspects of filler reinforcement were discussed by using specific test or literature results. They can be summarized as follows: under certain conditions, the polymer chains are adsorbed onto the clay filler surface or react chemically. In this case, when the polymer chains are immobilized at the vicinity of the filler surface, the modulus of this region is increased and acts more as a filler than as a polymer. Hence, the effective filler volume fraction is enhanced at the expense of the effective free polymer. Depending on the thickness of the glassier layer around the filler particles and the gradient of its modulus, the macroscopic properties can be influenced to a greater or lesser degree. Therefore, surface area, surface characteristics, and the structure of the filler influence the reinforcement effects. In this context, also the filler content plays an important role. It is generally assumed that the polymer–filler interaction can be promoted by the increase in the distance of the single silicate layers, for example, through cation exchange. Here, the type of clay modifier is very important as it is described in other chapters of this book. Furthermore, the intercalation of polymer chains is the first step to exfoliate the layered silicate into single platelets, which should be uniformly dispersed within the matrix. In most of the literature, it is stated that a complete exfoliation of the filler is a precondition for the highest reachable property level of the compounds as it was discussed in the previous sections. However, a combination of intercalated and exfoliated parts of the filler is often found in real materials. It definitely needs a certain degree of filler dispersion and polymer–filler interaction to achieve an effective mechanical reinforcement. In the following paragraphs, some selected aspects of clay reinforcement are discussed in more detail, but it can be said that particle size, polymer–filler interaction, intercalation/exfoliation state, and filler dispersion can hardly be discussed independent of each other. For rubber/clay materials, various compatibilizers have been described in the literature, such as epoxidized natural rubber (ENR), maleic anhydride (MA), silane, or resorcinol. Another possibility is to create master batches containing the filler in case the compatibility of the polymer and the filler is not high enough to disperse the filler well into the polymer matrix, for example due to different polarity. The principle aim of the application of compatibilizers is to achieve additional chemical bonding to increase the polymer–filler interaction because they can increase the miscibility of the polar filler and a non-polar polymer such as SBR, NR, or EPDM [22]. A better miscibility can contribute to a better dispersion of the filler, which is an important factor determining the physical properties of the resulting material. Another aspect is the possible increase in the interlayer distance of the plate-like filler through

330

MECHANICAL AND FRACTURE MECHANICS PROPERTIES

the molecules of the compatibilizer partly being of small size, which, in fact, supports the intercalation/exfoliation process [17,22,24,28]. For example, Di Gianni et al. [85] showed that the use of MA for rubber/OMMT materials results in a well-dispersed and intercalated filler morphology with comparably increased interlayer distances. Silane should react with the filler or should be adsorbed on the filler surface. It should furthermore have an affinity to the polymer [14]. In Ref. [86], the principle of the effect of the silane was described as follows: First, silane molecules are grafted onto the surface of the silicate platelets, leading to a decreasing surface energy of the clay and therefore to an increase in miscibility between the clay and the polymer, in this case polyethylene terephthalate (PET). This results in a better dispersion of the filler. According to Ref. [39], where EVA/MMT materials were investigated, there is a possible reaction between the silanol groups of the silane and the hydroxyl groups at the surface of the MMT. Regarding the aspect of functional groups at the filler surface, the kind of ion exchange together with the type of compatibilizer will play an important role. Concerning a better dispersion of the clay platelets, the more effective shearing during mixing in the case of an enhanced polymer–filler interaction due to the silane can result. Better clay dispersion in comparison to the unmodified material was found, for example, in PDMS materials modified with silane [38]. A study of Kim et al. [18] showed that the mechanical properties of NBR materials could be increased by using a certain type of silane. The authors explained their findings with a stronger chemical interaction including the formation of Si–O–Si bonds between the MMT and the silane, as well as Si–C bonds between the silane and the polymer. However, an optimum of the silane content was found at 5 phr. The reason for the decreasing strength was said to be the formation of defects in the nanocomposites caused by adding a large amount of silane. Also, the authors’ own results of instrumented tensile-impact tests and tensile tests with elastomeric compounds on the basis of NR/BR/SBR reinforced with two bentonite fillers showed an influence of the silane content depending on the type of the filler. For one of the fillers combined with a large amount of silane, a strong decrease in toughness and strength was obtained. These results indicate that silane can have a negative influence when it is overdosed. In Ref. [85], it was proposed that under certain conditions, a kind of elastic coating on the filler surface is obtained due to the use of a coupling agent, triggering chemical reactions on the filler surface. This coating may lead to the improvement in mechanical properties because during a mechanical loading, it may act as a kind of “damper.” Thus, the stress between polymer and filler is “softened” [87]. Also, Ramorino et al. [1], who investigated NR/MMT materials, found a decreasing material stiffness, expressed by the initial modulus from quasistatic tensile tests, with increasing test temperature for NR materials with higher clay contents. As a reason, the “melting” of the glassy interface between filler and polymer was given. Varghese et al. [11] suggested on the basis of DMA tests that three different states of clay filler appear in ENR: intercalated, exfoliated, and “confined”. The reason for this confinement (re-aggregation) is the chemical reaction between the intercalant and the sulfur-based cross-linking system. This means that the latter should be selected very carefully because of its possible effect on the intercalation/exfoliation behavior.

MECHANISM OF REINFORCEMENT IN RUBBER–CLAY COMPOSITES

331

As another example, the study in Ref. [41] should be mentioned. The experiments with EPDM/organoclay materials showed that the compatibilizer MA can increase the effective cross-link density of peroxide-cured EPDM/organoclay materials. In addition, the dispersion of the clay was better with the compatibilizer and full exfoliation was reached. Another possibility to increase the compatibility of rubber and filler was shown by Chakraborty et al. [28], who investigated SBR/bentonite vulcanizates with various amounts of filler. The aim of their study was to characterize the influence of a resol-resin modification of the filler on the material properties. Some results are summarized in Figure 10.12. The authors found a larger interlayer distance of the clay when it was modified. This means that at least an intercalation took place. Additional TEM investigations showed that an intercalated/exfoliated structure appeared. From these results, it can be derived that the clay content has an optimum value at 10 phr. When using more than 10 phr clay, both tensile and tear strength decrease. It was interesting to find that the portion of bound rubber (BR) and the difference of the maximum and minimum torque Dtorque from vulcameter tests which can be taken as a measure of the cross-link density have also a maximum at 10 phr. Because the portion of bound rubber is a measure of the polymer immobilized at the filler surface, the results presented in Figure 10.12 indicate a reduction in the polymer–filler interaction at a filler content of 15 phr, which may be due to filler agglomeration. Beside the polymer–filler interaction, the primary particle size of the used filler plays an important role for the reinforcement effect of fillers. Smaller particles can be uniformly dispersed within the polymer matrix more easily [87]. Furthermore, the contact area of fillers with smaller particles is larger than that of fillers with greater particles. This is in direct conjunction with the resulting polymer–filler interface. Deng et al. [2] used three types of non-specified “red clay” as fillers for NR 50

10 Δ torque

Bound rubber σmax

40

8

30

6

20

4

10

2

0 0

2

4

6

8

10

12

14

Δtorque (dNm)

Ts (N/mm); σmax (MPa); BR (%)

Ts

0 16

Clay content (phr)

FIGURE 10.12 Tear strength Ts, tensile strength smax, bound rubber, and the difference of the maximum and minimum torque Dtorque from vulcameter tests as a function of the content of the modified clay; data from Ref. [28].

332

MECHANICAL AND FRACTURE MECHANICS PROPERTIES

vulcanizates. For each of the clays, four different primary particle sizes were realized ranging from 0.07 to 1.4 mm. It could be observed that the mechanical parameters of the elastomers were partly strongly affected by the particle size of the filler. The largest particle size led to the lowest tensile strength, tear strength, and wear resistance, while the smallest particle size resulted in a balanced property level. The cross-link density of the material may be influenced by addition of organoclay [40]. The reason is the adsorption or reaction of the cross-linking agent(s) on/with the filler surface [29,43]. Apart from this, the cross-linking system of rubber–clay composites is also important regarding the resulting mechanical properties. In the literature, different mechanisms have been described that are of chemical nature and may lead to an increase in the distance between the silicate layers and therefore to an increased free surface area of the filler. In Ref. [16], it was stated that by using a cross-linking system based on sulfur as well as a peroxidic one, an intercalated/exfoliated silicate structure in acrylonitrile rubber (NBR) could be achieved. However, the orientation of the silicate layers was higher when the crosslinking was performed with a peroxidic cross-linking agent. Furthermore, it was found that the level of exfoliation in the sulfur-cured materials became higher with increasing content of stearic acid. This was traced back to the better dispersion of the filler through the stearic acid, whose long molecules move into the galleries of the layered silicate, increase the gap distance, and promote the exfoliation process [15,16]. For chloroprene rubber, it was found that the use of unmodified MMT and layered double hydroxide (LDH) led to a decreased maximum rheometric torque in comparison to that of the mixtures with the organically modified versions of these fillers [36]. As a reason for this behavior, the inactivation of the cross-linking agent was discussed. Furthermore, not only the result of the cross-linking reaction but also its kinetics is influenced by the nature of the clay. This effect was discussed in detail in terms of chemical reactions between the clay surface and the cross-linking agents and is not consistent for the different types of clays investigated. Another example is given in Ref. [37], where investigations with polydimethylsiloxane rubber–clay composites with two different kinds of MMT (natural and organically modified) were described. The authors found that unmodified MMT led to stronger polymer–filler interaction compared to the organoclay because of covalent bonds between the polymer chains and the filler surface, which contributed to the overall cross-link density. In contrast, the polarity of the OMMT surface led to decreased polymer–filler interactions. Furthermore, the cross-link density of the rubber/OMMT materials was lower and decreased with the filler content because of a chemical reaction of the filler and the cross-linking agent at the filler surface, leading to a reduction in the effectively available number of radicals for the rubber cross-linking reaction. A similar effect for PDMS/clay compounds was described in Ref. [38]. A further effect contributing to an enhanced level of physical properties is the ability of some rubbers to crystallize under strain. Here, natural rubber and chloroprene rubber are some examples. The crystallites developing under strain act as additional cross-linking points or as a kind of filler, which leads to a selfreinforcement of such elastomers. As a consequence, for example, the tensile strength is increased by this effect. Results of tensile tests with NR/clay compounds [6]

THEORIES AND MODELING OF REINFORCEMENT

333

showed that the tensile stress in the beginning of the stress–strain curve is larger for NR/clay materials, but the characteristic shape of a strain-crystallizing material is missing for clay contents larger than 5 phr. Therefore, the authors assumed hindered strain crystallization because of the presence of the clay. A contrary result was obtained by Qu et al. [12], who reported an increase in tensile behavior of NR vulcanizates especially with smaller amounts of OMMT. In situ wide-angle X-ray diffraction (WAXD) tests showed the most pronounced increase in crystallinity with increasing deformation for the material with 5 phr OMMT, for which the highest tensile strength was determined. Furthermore, the onset of crystallization and the resulting crystallite size were reduced by addition of the filler. Similar observations were made by Carretero-Gonzalez et al. [88], who found additionally a two-step regime of the strain-induced crystallization of NR with 15 phr OMMT. As the structural reason, the authors stated the increased interfacial interaction between filler and polymer, as well as an enhanced mobility of the polymer chains at the filler surface, as factors promoting the strain-induced crystallization. Furthermore, in strong relation to these findings is the exfoliation of the filler because of the larger surface area. Because this all is influenced by the surfactant of the MMT, the MMT must have an influence also on the crystallization behavior. In Ref. [36], results were reported indicating that OMMT promotes strongly the strain-induced crystallization of CR vulcanizates, while for MMT such a pronounced effect was not found. In Ref. [2], the development of a (semi)-rigid clay network within the matrix due to more or less strong hydrogen bonding between the particles is given as a possible mechanism of reinforcement. The level of bonding depends on the number of hydroxyl groups. Ramorino et al. [1] showed that the percolation threshold for the filler networking is given for the NR/OMMT materials at a clay content of 8 phr, connected with a more pronounced stiffness of the materials. In conclusion, it can be stated that the mechanisms of reinforcement in rubber/clay materials are very complex. It is difficult to discuss these mechanisms regarding the influence of the different aspects on the mechanical properties in general terms because of the great number of factors influencing the morphology development. Various aspects such as filler modification, possible chemical reactions taking place between the mixture ingredients, especially the cross-linking system and the filler surface, resulting intercalation and exfoliation of the filler, which are responsible for a good dispersion of the filler, and finally interactions between the polymer and the filler must be taken into account when discussing the resulting properties.

10.5 THEORIES AND MODELING OF REINFORCEMENT Modeling the material behavior may provide valuable information on the underlying physics and contribute to the establishment of structure–property relationships [89]. During the past years, many efforts have been made to develop and apply numerical simulation methods for the prediction of certain properties of filler-reinforced polymers including elastomers ranging from molecular simulations to micromechanical models. The analytical models for the prediction of the materials’ stiffness can

334

MECHANICAL AND FRACTURE MECHANICS PROPERTIES

be classified into two main groups: bounding models and mesomechanics models [91]. General information about molecular simulations, theories, and different models can be found in Refs [48,90], while Ref. [91] gives a short overview. In general, the models for the prediction of composite properties assume well-dispersed rigid particles in a much softer matrix. This means there are two domains, the particle and the matrix domain, which are treated as homogeneous material with certain elastic properties [92]. Various continuum-based micromechanical models for the prediction of mechanical properties – and here especially the modulus as a measure of the stiffness – have been developed or tried to be transferred to composites with plate-like fillers, for example, according to Guth, Halpin–Tsai, Mori and Tanaka, and Hui and Shia [1,84,89,93–99]. In principle, a polymer–clay nanocomposite can be considered as a set of geometric (particle size, aspect ratio, etc.) and “material” features such as exfoliation/intercalation state or polymer–clay interface behavior [92]. When modeling the macroscopic behavior of polymeric composites, generally a homogeneous distribution of the filler particles and a defined particle–matrix structure is assumed. If filler particle aggregation and/or intercalated structures occur, as it is often the case for polymeric nanocomposites, the predictions may fail. Furthermore, for the prediction of the modulus of nanocomposites, some assumptions are made, which are inherent with the reality [89,92]: . . .

Filler and matrix behave linearly elastic, are isotropic and perfectly bonded. Filler particles are perfectly aligned, asymmetric and with uniform particle shape and size. Polymer–filler interactions are not considered.

Another precondition for robust predictions is the application of a reliable method for characterization of the intrinsic morphology of the intercalated clay and of the matrix morphology and properties near the filler particles [92]. For these reasons, the results of micromechanical modeling are not always in good agreement with experiments and to solve such problems, among others, the structural parameters of the nanofiller should be carefully handled [89] and generally further development is necessary. The following equation (Eq. (10.9)) is known as the Guth equation [93], which is applicable to the prediction of the modulus of elasticity for polymers with spherical particles and lower filler contents up to 10 vol.%: E ¼ Em ð1 þ 2:5f þ 14:1f2 Þ

ð10:9Þ

with E denoting the modulus of elasticity of the filled polymer, Em the modulus of elasticity of the polymeric matrix, and f the volume fraction of the filler. Furthermore, Eq. (10.10) was amended by a shape factor a to extend the applicability to non-spherical filler particles with a ratio between length and width  1: E ¼ Em ½1 þ 0:67af þ 1:62ðafÞ2 

ð10:10Þ

THEORIES AND MODELING OF REINFORCEMENT

The Halpin–Tsai equations are generally of the following form:   ð1 þ jhfÞ E ¼ Em ð1hfÞ

335

ð10:11Þ

with j denoting the shape factor, and h ¼ ðEf =Em 1Þ=ðEf =Em þ jÞ with Ef denoting the modulus of elasticity of the filler [1,95]. Wu et al. [95] found that the predicted values of E according to Eq. (10.11) are higher than the experimental ones. By introduction of a modulus reduction factor (MRF) of 0.66, the predicted values fit very well to the experimental results also with higher clay contents. According to the authors, this MRF is necessary because the morphological characteristic of the plate-like filler is disregarded in the theory. As a result, not only the Guth equation but also the modified Halpin–Tsai equation could predict the modulus of SBR, NBR, and CNBR materials with layered silicates very well. Similar results were obtained by Gatos and Karger-Kocsis [96], who modified the Guth equation for prediction of the modulus of HNBR/layered silicate materials. It could be shown that in the whole range of filler content investigated, the modulus could be predicted very well by using Guth equation with the modification with MRF according to [95]. Also, Frogley and coworker [98] found a reasonable agreement of the experimental data and the prediction with the Guth equation for single-wall nanotubes and carbon nanofibers with aspect ratios of 70 and 120. However, with a much larger aspect ratio of 1000, prediction of the modulus failed. According to the authors this is due to the fact that the increase in the ratio between surface area and volume is ignored. The authors [93] extended the above equations by a parameter called interface area function (IAF) and could therefore well predict the modulus of elastomeric composites with layered silicate. Results of experiments and predictions by using Guth and Halpin–Tsai equations shown in Ref. [97] are not in perfect agreement. Especially with medium contents of layered silicate, there are strong deviations between the experimental and the theoretical results. This was traced back to the fact that the materials contained also intercalated structures, which were, however, disregarded in the models. Beside the above-described models, other possibilities exist. For example, Buxton and Balazs [100] used the lattice spring model to compare the reinforcement effects of plate-like and spherical filler particles also in a viscoelastic matrix. They found that plate-like fillers reveal a much more pronounced reinforcement effect than particulate fillers. Furthermore, it was stated that the reinforcement effect is larger when the clay is fully exfoliated. If the clay is in an intercalated state, the stress concentrations in the vicinity of the clay stacks are lower compared to those at single platelets. Because the reinforcement level depends on the distribution of local deformations, in the case of exfoliated morphology a higher reinforcement is therefore to expect. A similar finding was reported by Yung et al. [89], who found an increasing modulus for epoxy–MMT composites with increasing degree of exfoliation by the application of the Halpin–Tsai model. However, “reinforcement” is often a synonym only of the increase in the composite stiffness. The “multiple” stress concentration fields around a stack of

336

MECHANICAL AND FRACTURE MECHANICS PROPERTIES

platelets with a certain interlayer distance, shown in Ref. [100], may be a cause for a better strength or toughness of materials with intercalated structure as it was described above. In conclusion, it can be said that application of well-known models for prediction of the elastic modulus of elastomeric composites with layered silicate is possible and leads to mostly good agreements between experimental and theoretical results, when factors are introduced, which regard the special geometrical characteristics of the layered silicates and the resulting in-rubber morphology. The additional combination with models describing the polymer–filler interaction such as the Kraus model [49] may lead to better understanding of the material behavior of polymeric materials reinforced with layered silicates. ACKNOWLEDGMENTS The authors would like to thank the German Research Foundation (DFG) for the financial support of the subproject 2 “Technical Material Diagnostics – Fracture Mechanics of Filled Elastomeric Blends” within the framework of the Research Unit FOR 597. Furthermore, we acknowledge the Continental Reifen Deutschland GmbH (formerly Continental AG), Hannover, Germany and Dr. M. Galimberti from Pirelli Tyre Company, Milano, Italy for providing elastomer materials. Dr. S. Ilisch from the Martin-Luther-University Halle-Wittenberg, Germany is acknowledged for performing DMA tests. REFERENCES 1. Ramorino, G.; Bignotti, F.; Pandini, S.; Ricco`, T. Mechanical reinforcement in natural rubber/organoclay nanocomposites. Compos. Sci. Technol., 69, 1206–1211 (2009). 2. Deng, H.; Ao, N.; Yan, Z.; Zhang, L.; Huang, C. The reinforcement of red clay on natural rubber and its reinforcing mechanism. J. Appl. Polym. Sci., 112, 3418–3422 (2009). 3. Lopez-Manchado, M. A.; Herrero, B.; Arroyo, M. Organoclay–natural rubber nanocomposites synthesized by mechanical and solution mixing methods. Polym. Int., 53, 1766–1772 (2004). 4. Sharif, J.; Yunus, W.; Md Z. W.; Mohd Dahlan, K. Z. H.; Ahmad, M. H. Preparation and properties of radiation crosslinked natural rubber/clay nanocomposites. Polym. Test., 24, 211–217 (2005). 5. Jurkowska, B.; Jurkowski, B.; Oczkowski, M.; Pesetskii, S. S.; Koval, V.; Olkhov, Y. A. Properties of montmorillonite-containing natural rubber. J. Appl. Polym. Sci., 106, 360–371 (2007). 6. Wang, Y.; Zhang, H.; Wu, Y.; Yang, J.; Zhang, L. Structure and properties of straininduced crystallization rubber–clay nanocomposites by co-coagulating the rubber latex and clay aqueous suspension. J. Appl. Polym. Sci., 96, 318–323 (2005). 7. Ramorino, G.; Bignotti, F.; Conzatti, L.; Ricco`, T. Dynamic and viscoelastic behavior of natural rubber/layered silicate nanocomposites obtained by melt blending. Polym. Eng. Sci., 47, 1650–1657 (2007).

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CHAPTER 11

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY BRENDAN RODGERS WEIQING WENG JOHN SOISSON DAVE LOHSE WALTER WADDELL ROBERT WEBB

11.1 INTRODUCTION Elastomeric nanocomposites used for air barrier applications consist of a base rubber and a silicate. The silicate or clay is typically treated with an organic surfactant such as a quaternary ammonium salt to aid compatibility. The inorganic component of the nanocomposite can be as much as 10 wt.% of the composition before compounding. Compared to the base elastomer, such nanocomposites are reported to have higher tensile modulus, storage modulus, improved heat resistance and thermal stability, and lower permeability. In industrial applications, low permeability has emerged as the most important benefit. Use of nanocomposites in tire innerliners with butyl rubbers as the base elastomer results in better tire air retention qualities and consequently, potential performance benefits. 11.1.1 Butyl Rubbers as Nanocomposite Base Elastomers Butyl rubber is a copolymer of isobutylene and approximately 2 mol% of isoprene [1]. The length of the isobutylene structural unit is 0.270 nm and is thus 67% of that of the 1,4-isoprene structural unit at 0.405 nm [2]. Isoprene is incorporated in a trans-1,4 Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc. Chapter 11 “Permeability of Rubber Compositions Containing Clay” by Michael Brendan Rogers
2011 ExxonMobil Chemical Company, a division of Exxon Mobil Corporation. Published 2011 by John Wiley & Sons, Inc.

343

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PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

enchained head-to-tail arrangement to produce a random, linear copolymer. Isobutylene polymerizes in a head-to-tail sequence producing a rubber having no asymmetric carbon atoms. The geminal-dimethyl group has two methyl groups bonded to the same carbon atom [–C(CH3)2)–CH2–] on alternative chain atoms along the polyisobutylene backbone producing a steric crowding effect. Distorting the hydrogen atoms of the methylene carbon (–CH2–) from the normal tetrahedral 109.5–124
, and the dihedral angle of the carbon–carbon single bond backbone by about 25
relieves some strain. The stereochemistry of the isobutylene unit thus results in close packing along the polymer chain, low free volume fraction, and consequently low permeability. Butyl and halobutyl rubbers are therefore a preferred platform for nanocomposite compositions when lower permeability is required. Currently halogenated butyl rubber, that is, chlorobutyl rubber and bromobutyl, are used when low permeability is needed. Halogenation of the butyl rubber improves vulcanizate properties. The halogen content of chlorobutyl or bromobutyl rubber is typically expressed in weight percent (wt.%). Figure 11.1 shows the structures of the bromine containing functional groups found in bromobutyl rubber [3]. For bromobutyl rubber the total bromine content includes both organic bromine, as shown in Figure 11.1, and inorganic bromine such as calcium bromide. Structure II, the predominant structure in bromobutyl rubber, represents 50–60%, followed by Structure I representing from 30 to 40%. Approximately 5–15% is Structure III. Structure IV is typically only 1–3%. CH3 CH 2

C

CH3 CH2

C

CH

CH2

(0.8–3.0 mol%) CH3 Bromination

CH2

CH3 CH2

C

CH

CH2

CH2

CH

C

CH2

Br Structure I

Structure II

CH3

CH2Br CH2

C

CH

CH2

CH

C

CH

CH2

Br Structure III

Structure IV

FIGURE 11.1 Structure of isoprenyl units in Bromination rubber [1,3].

INTRODUCTION

345

Similar structures are found in chlorobutyl rubber but the relative ratios of the four respective configurations changes. For example, no evidence has been found for Structure III in chlorobutyl rubber [1]. Tire innerliners use predominantly bromobutyl rubber. Due to the lower carbon–bromine bond energy rendering it more reactive, bromobutyl rubber innerliners show better liner to tire casing adhesion and cure rate compatibility with adjacent tire components when compared to regular butyl or chlorobutyl rubber. Permeation of a gas or permeability through a membrane such as a tire innerliner consists of three distinct processes [4,5]. First, the gas molecules must dissolve on one side of the membrane, then diffuse across the membrane or innerliner to the opposite side of lower gas concentration, the rate being dependent on the size of the diffusion gradient, and then evaporate or disperse in the adjacent tire component or other medium. The rate of diffusion of oxygen and nitrogen through a butyl, chlorobutyl, or bromobutyl rubber membrane is a function of a number of parameters. From van der Waal’s interactions the size of
oxygen and nitrogen molecules are approximately 2.9 and 3.1 A, respectively, with the smaller molecule expecting to diffuse more rapidly. This is observed in practice. Amerongen, in calculating the energy of activation of gas diffusion, also noted that diffusion coefficients could increase between 1.2 and 3 times with each 10
C increase in temperature. The temperature dependence can be described using the expression: D ¼ Do expð Ed =RTÞ

ð11:1Þ

Here, D is the diffusion coefficient, Do is a frequency factor, and Ed the activation energy of diffusion. For regular butyl rubber used in, for example, a tire curing bladders, the solubility of oxygen and nitrogen are 0.155 and 0.122 cc gas/cc of polymer, respectively, in accord with oxygen’s faster diffusion rate. The solubility of a gas is also a function of temperature. As temperature increases, the diffusion rate of oxygen also increases [6]. The activation energy of diffusion was reported to be 50.6 and 49.8 kJ/mol, respectively, for nitrogen and oxygen with frequency factors, analogous to the Arrhenius constant of 43 and 34 cm2/s. Thus, from the thermodynamics this suggests that oxygen diffuses more readily than nitrogen, as is observed [5]. 11.1.2 Measurement of Tire Innerliner Compound Permeability A method for measurement of oxygen or other gas permeability through a vulcanized rubber sheet is by use of a Mocon Ox-Tran Model 2/61 oxygen transmission rate test apparatus and Perm-Net operating system [7]. In this specific instrument, there are six cells where gas transmission through each test sample placed in a cell is individually measured. A zero reading to establish a baseline is obtained and samples can be tested typically from room temperature up to 60
C. Data is reported as a permeation coefficient in cc mm/(m2-day), that is, transmission of a volume of gas through a 1 mm cross section, per square meter surface area of test specimen, per unit time (day), and permeability coefficient in cc mm/(m2-day-mmHg) that is the permeation

346

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

TABLE 11.1 Model Bromobutyl Screening Compound [8] Material Bromobutyl rubber Carbon black (N660) Naphthenic oil Aromatic hydrocarbon resin Phenolic tackifying resin Stearic acid Zinc oxide Mercaptobenzothiazole disulfide (MBTS) Sulfur

phr 100.00 60.00 8.00 7.00 4.00 1.00 1.00 1.25 0.50

Note: Bromobutyl rubber used was EXXON bromobutyl 2222 and EXXON bromobutyl 2255 [9].

coefficient corrected for atmospheric pressure. For tire innerliners, permeability can then be expressed as a rating relative to a control compound. The model compound illustrated in Table 11.1 is assigned a rating of 100. This control compound has a nominal permeation coefficient at 60
C of 500–550 cc mm/(m2-day). Additional information on the test method can be obtained in ASTM D3985. 11.1.3 Further Improvement in Tire Permeability The least permeable elastomer compounds used in tire innerliners consists of a carbon black filled bromobutyl or chlorobutyl rubber compound with only sufficient process oil to meet product manufacturing needs. Isobutylene-based nanocomposites can offer an option to significantly reduce permeability. Such systems can use a modified butyl rubber, brominated isobutylene-co-para-methylstyrene (BIMSM) and an organoclay. 11.2 NANOCOMPOSITES Elastomer nanocomposites can have a number of benefits over the base polymer such as increased stiffness, improved impact resistance, flame retardance, and permeability. Of this range of properties gas barrier capability is of particular interest. Significant reductions in permeability of rubber compounds containing nanocomposites and specifically nanoclays have been reported and can offer performance benefits such as improvement in tire inflation pressure retention (IPR) [10,11]. Elastomer nanocomposites essentially consist of the base polymer and a nanoclay that is typically a montmorillonite. The montmorillonite clay may contain a modifier that aids in the dispersion of the clay within the polymer matrix and where the clay can adapt one of the five states. Particulate Dispersion. The organoclay particle size is in the order of microns but uniformly dispersed. The terms, aggregate, and agglomerate, have frequently been used to describe this state.

NANOCOMPOSITES

347

1. Intercalated Nanocomposite. This describes insertion of polymer chains into the layered silicate structure and occurs in a crystallographic regular fashion, regardless of the polymer to clay ratio. Intercalated nanocomposites may typically contain several layers of polymer between organoclay plates. Increase in the gallery spacing of a nanoclay, swollen in rubber, from a pristine state of 0.3–0.7 nm up toward 2.0–6.0 nm can be considered as creating an intercalated condition. 2. Flocculated Nanocomposites. Conceptually this is the same as intercalated nanocomposites. However, the silicate layers are sometimes flocculated or aggregated due to hydroxylated edge-to-edge interactions of the silicate layers. 3. Intercalated–Flocculated Nanocomposite. Though the clay plates can be separated, tactoids or agglomerates can be formed that are in the order of 100–500 nm. 4. Exfoliated Nanocomposites. In an exfoliated nanocomposite the individual clay layers are separated within a continuous polymer matrix, by an average distance that depends on the clay concentration or loading. The clay plates in the nanocomposite essentially function by increasing the path length that gas molecules must travel. The term “tortuous path” has been used to describe the consequent retarded flow rate of gas molecules through a nanocomposite shown in Figure 11.2 [10–12]. The dimensions and orientation of the nanoparticle plates are important in defining the degree of tortuosity in that when, for example, nanoclay particles are parallel to direction of gas flow their impact will be negligible. Conversely, when the plates are perpendicular, they will arrest or hinder gas molecules flowing through the nanocomposite. Models can be developed to describe the effect of orientation. In an idealized system with clay plates perpendicular to the gas flow, as clay plate diameter or aspect ratio increases permeability is reduced, and as the clay volume fraction increases, again permeability decreases. No clay

Clay

Direction of flow

W (width) L (length)

FIGURE 11.2 Illustration of tortuous path due to introduction of nanoclay.

348

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

Compared to spheres or fibers, the plate-based morphology is most effective at increasing the path length due to the large length-to-width ratio. A tortuosity factor, T, can be described as the ratio of the actual distance, d00 , that a gas molecule must travel to the shortest distance, d0 , that it would travel in the absence of clay plates dispersed in the base polymer. It can be expressed in terms of the length L; width W; and volume fraction of the clay plates, f, as 00

T ¼ d =d 0 ¼ 1 þ ðL=2WÞf

ð11:2Þ

The tortuosity effect on the permeance or permeability coefficient could then be expressed as Pn =Pp ¼ ð1  fÞ=T

ð11:3Þ

where Pn and Pp represent the permeability of the nanocomposite and pure polymer, respectively, f is the inorganic clay volume fraction, and T is the tortuosity factor. For montmorillonite, the polygonal aluminosilicate sheets can be approximated as discs with a mean diameter L ranging from 50 to 120 nm and a nominal width of 1 nm, excluding the gallery spacing. This calculation, sometimes referred to as the Nielsen’s model, has been reported to work well in developing empirical estimates of gas barrier properties illustrated in Figure 11.3 [13]. However, gas permeability measurements in many instances are significantly above the theoretical prediction, especially in elastomer nanocomposites. A key assumption of this description is that the sheets are placed in an arrangement such that the direction of diffusion is perpendicular to the direction of the sheets. This arrangement results in the highest tortuosity, and any deviation from the arrangement where the sheet orientation deviates from the film plane would lead to a shift of the barrier properties. This preferred condition is difficult to achieve in practice due to reaggregation of the clay plates and association of the plates with other materials in a compounded nanocomposite.

Relative gas permeabilty

1.2 1

1 2 20 50 100 200 300 1000

0.8 0.6 0.4 0.2 0 0

0.02

0.04

0.06

0.08

0.1

0.12

Volume fraction

FIGURE 11.3 Theoretical gas permeability calculation (silicate plates perpendicular to the gas flow). Effect of aspect ratio and clay concentration (volume fraction).

NANOCOMPOSITES

S = –1/2

S=0

349

S=1

FIGURE 11.4 Clay plate orientation.

Theoretical models generally assume that the silicate layers are aligned perpendicular to the direction of gas flow, are exfoliated, and display a large aspect ratio. However, permeability is a function of diffusivity (i.e., the diffusion coefficient) and equilibrium absorption of the gas. Clay silicate plate orientation, aggregate or tactoid formation, and uniformity of dispersion will also have an effect on permeability. The effect of clay plate or sheet orientation can be illustrated in Figure 11.4 [14,15]. In an idealized model, should the plates be orientated parallel to the flow of gas molecules, no effect on permeability might be expected. This would be in contrast to when the plates are perpendicular to the flow of gas or are in a random orientation. An “order parameter,” S, can be determined that would further quantify the relationship between the tortuosity factor and clay plate orientation, where S is S ¼ 1=2ð3 cos2 u  1Þ

ð11:4Þ

The term, S, represents the angle between the desired orientation and the direction of actual orientation in the polymer matrix. In this case when u ¼ 0, S ¼ 1 indicating orientation perpendicular to gas molecule flow and when S ¼  1/2(u ¼ 90
) orientation is parallel to the flow of gas. If S ¼ 0 sheet orientation is random. Equation (11.3) can then be modified as follows Ps =Pp ¼ ð1  wÞ=½1 þ ðL=2WÞws ð2=3ÞðS þ 1=2Þ

ð11:5Þ

This reduces to Eq. (11.2) when S ¼ 1, that is, when the sheets are perpendicular to the flow of gas, and converges to the permeability of the pure rubber when S ¼  1/2. In industrial application such as in tire innerliners, a number of factors will impact the degree of orientation of the clay plates such as the degree of processing, disorientation due to obstructions such as carbon black particles and loss of the surfactant on the clay surface due to excessive heat destabilizing the organoclay. An orientation parameter, S, in the range 0.20–0.35 can represent a reasonably high degree of orientation. In practice, the orientation parameter can be determined by small angle X-ray scattering (SAXS). X-ray measurements for nanocomposite samples are taken by orienting the sample so that the beam hits it at two different angles with respect to the plane of the film, that is, in the edge-on and face-on directions, respectively [8,10].

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

350

FIGURE 11.5 SAXS scattering patterns of BIMSM montmorillonite compound showing clays are oriented along the flow direction film plane [10].

Reactor-extruded sample

1

Edge-on Face-on

0.1

Compound sheet

10

Edge-on Face-on

1

Intensity [a.u.]

Intensity [a.u.]

10

Intensity [a.u.]

In the SAXS profile of an organoclay, scattering peak positions are typically nonequidistant, indicating multimodal distributions of thickness of organomineral layers, where layers overloaded by modifier coexist with layers having low surfactant content within the same stacks or tactoids. From SAXS profiles of elastomer nanocomposites, when the organoclay peak due to the layer-stacking period almost disappears, this can suggest that the organoclay stacks are partially exfoliated in the polymer matrix. As an example, SAXS data was collected for two different sample orientations. The corresponding SAXS patterns collected along two sample orientations (face-on and edge-on view) confirmed that intercalated clays were oriented parallel to the film plane (Figure 11.5). Orientation of the clay plates can be improved during the compounding and tire building process. As shown in Figure 11.6, the clays in reactor product have little orientation. However, a strong orientation could be developed in the innerliner compound as it is processed into a sheet and then stretched during the tire building process. Clay plate orientation may be further quantified. Assuming organoclay stacks have a cylindrical symmetry, the orientation of organoclays can be described by an orientation distribution function g(b) depending on a single angle b that is defined as the angle between the normal of organoclay stack and the normal of the film plane.

Edge1 Edge2 Face Innerliner in tire

10

1

0.1 0.2

0.4 0.6 S [nm–1]

0.8

0.2

0.4 0.6 S [nm–1]

0.8

1

2

4

3

5

6

2θ

FIGURE 11.6 SAXS profiles of BIMSM montmorillonite samples at different process stages, showing the development of clay orientation in the process [10].

NANOCOMPOSITES

351

Azimuthal sections for the SAXS analysis of lamellar stacks was carried using a modified Onsager orientation distribution function: gðbÞ ¼ P0 þ ð1  P0 Þ*ðP=sin hPÞ*ðcoshðP cos bÞÞ

ð11:6Þ

The extent of the preferred orientation of these organoclay stacks can be calculated using Herman’s orientation parameter:  2 ¼ 0:5*½3* P

p=2 ð

gðbÞ cos2 ðbÞsinðbÞd b  1

ð11:7Þ

0

A relationship can be established between orientation of the organoclay with the weight percent of organoclay added as a nanofiller dispersed in polymer matrix. Three-dimensional modeling of clay stacks in nanocomposites can be carried out and the extent of the orientation of these organoclay stacks calculated using the Herman’s orientation parameter:  2 ¼ ð1  P0 Þð1  3p  1 ½cothðpÞ  p  1 Þ P

ð11:8Þ

The Onsager orientation distribution function as defined in Eq. (11.6) has been used for SAXS analysis of lamellar stacks in presence of carbon black and Herman’s orientation parameter. Equation (11.8) is used to calculate the orientation of organoclays in polymer matrix. The calculated Herman’s orientation parameters for compounded BIMSM montmorillonite samples are compared in Table 11.2. The orientation seems to increase a little upon curing. Interestingly, the d-spacing between clay platelets also decreased a little upon curing. This may be the result of tight polymer network structure after curing, which compressed the layered clay structure. The d-spacing of 4.2–4.4 nm is significantly greater the organoclay montmorillonite (d-spacing 2.4 nm), indicating that even within the unexfoliated clays, they are intercalated by the polymer molecules. Table 11.3 shows the orientation factor values of the reactor produced nanocomposite sample and the innerliner specimen cut from a cured tire. As discussed previously, a strong orientation is developed in the tire building process. When comparing the orientation factors between the edge-on and face-on direction, the larger difference in the innerliner sample the greater the clay plate alignment. TABLE 11.2 Calculated Herman’s Orientation Parameters and d-Spacing for Compounded BIMSM–Montmorillonite Nanocomposites [15] Sample Nanocomposite compound (uncured) Nanocomposite compound (cured) Nanocomposite compound with DBTU cure system

Orientation

d-spacing (nm)

0.56 0.57 0.60

4.44 4.29 4.29

352

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

TABLE 11.3 Calculated Herman’s Orientation Parameters and d-Spacing for Compounded BIMSM–Montmorillonite Nanocomposites [15] Sample Reactor-extruded Innerliner in the tire

Orientation (face-on)

Orientation (edge-on)

0.2369 –

0.4622 0.3069

11.3 PREPARATION OF ELASTOMER NANOCOMPOSITES There are three primary methods for producing rubber–clay nanocomposites [8,11]. 1. Melt mixing where the elastomer and nanoclay are added together in an internal mixer with other compounding materials. 2. In situ polymerization where the nanoclay is added to the monomer in solution or in an emulsion, and polymerization can occur within the clay plates. 3. Solution blending where the nanoclay, in a suspension, is added to the rubber while in a solution or cement. Coagulation of polymer latex in the presence of aqueous clay suspension could also be used. A comparative study of melt versus solution preparation of a butyl–nanoclay composite was conducted by Liang et al. [16]. The authors found 1. the aspect ratio of clay layers in a melt mixed butyl nanocomposite was smaller than that for a solution prepared nanocomposite; 2. from hardness data (Shore A), nanoclay layers have much greater reinforcing effects than micro layers (10  6 m) typically found with kaolin clays; 3. nanodispersed clay with high aspect ratio has greater stress bearing capability; 4. there is stronger interaction between clay layers and rubber molecules when the clay plates are larger (higher aspect ratio) due to the larger contact surfaces restricting motion of rubber chains; 5. gas barrier properties, measured as nitrogen permeability at 40
C, are better with solution prepared nanocomposites. The authors demonstrated that addition of clay to butyl rubber reduced permeability by over 20% and that solution blending was superior to melt mixing. A benefit in the order of 5–10% was achieved via solution blending.

11.4 TEMPERATURE AND COMPOUND PERMEABILITY To demonstrate the practical effect of temperature on nanocomposite compound permeability, two formulations were prepared and then tested for permeability [6] and discussed briefly:

TEMPERATURE AND COMPOUND PERMEABILITY

353

1. Compound 1 is a conventional bromobutyl-based industrial innerliner formulation. 2. Compound 2 is a nanocomposite-based formulation under development for tire innerliner applications. This is based on brominated isobutylene-co-paramethylstyrene (BIMSM) and a nanoclay. The nanocomposite is considered as 100 rubber hydrocarbon (RHC) with the total RHC set at 100 [7]. The formulation composition and associated fundamental properties of the two compounds are tabulated in Table 11.4. Permeation coefficients are shown in Table 11.5. TABLE 11.4 Model Bromobutyl and Nanocomposite Innerliner Formulations and Properties [6] Compound

1

BIIR 2255 Nanocomposite N660 Naphthenic oil Struktol 40MS Phenolic Tackifying Resin Stearic acid ZnO MBTS Sulfur

2

100.00

Mooney viscosity ML1 þ 4 (MU) Cure rheometer (160
C) DTorque (MH-ML) (dN m) t10 (min) t90 (min) Vulcanization rate (dN m/min) Tensile strength (MPa) Elongation (%) 300% modulus (MPa) Hardness A (Shore A) Tear strength (kN/m)

60.00 8.00 7.00 4.00 1.00 1.00 1.25 0.50

100.00 55.00 2.00 3.00 3.00 1.00 2.00 1.00 1.00

64.5

60.9

3.4 1.5 7.6

4.7 2.8 23.0

1.0 10.6 789.1 3.5 45.9 56.9

0.3 11.3 711.3 5.8 61.6 59.2

TABLE 11.5 Effect of Test Temperature on Permeation Coefficient [6] Compound Test temperature (
C) cc mm/(m2-day) cc mm/(m2-day) cc mm/(m2-day)

30 40 60

1

2

118 213 572

43 121 355

354

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

Increase in test temperature leads to a large increase in permeation coefficient. This increase appears to follow a simple exponential of the form, y ¼ ax, though more data is obviously needed to quantify the relationship. The rate of increase for the nanocomposite compound is less than that for the bromobutyl rubber compound, with the differential increasing as the test temperature increases. This would suggest that at high tire operating temperatures, where oxygen permeation into the tire casing is greater (i.e., increase in tire intracarcass pressure, ICP), nanocomposites will be more efficient and in turn, show better end product performance. This is in fact what was reported in earlier work using BIMSM-based nanocomposites in truck tires [10]. To further illustrate the effect of temperature on permeability, three model compounds were aged in an air oven under two sets of conditions, 7 days (168 h) at 100
C, and 3 days (72 h) at 125
C. The test samples consisted of the following 1. Bromobutyl rubber used in many industrial innerliner formulations. 2. Nanocomposite base polymer (brominated isobutylene-co-para-methylstyrene) in an identical formulation to that used for bromobutyl rubber. 3. A nanocomposite composition based on brominated isobutylene-co-paramethylstyrene and a montmorillonite nanoclay dispersed within the compounded rubber. For comparative purposes the formulation used was the same as that for bromobutyl rubber except the naphthenic oil content was decreased to compensate for other lower molecular weight additives used in the production of the nanocomposite formulation. The properties of the three compounds are shown in Table 11.6. A number of points were evident in the study: 1. Cure rate, t90, tensile strength, elongation, and 300% modulus are all within the range required for application in a tire innerliner. 2. With aging the permeability of a bromobutyl rubber-based innerliner will decrease, potentially due to the loss in low molecular weight components in the formulation. 3. Nanocomposite base polymer (BIMSM) formulations, identical to the bromobutyl rubber compound control, show significantly better permeability performance. The rate of decrease in permeability is also less suggesting a more stable formulation. 4. The nanocomposite based formulation showed the lowest permeability and the least amount of change with aging, suggesting it is a more stable system. BIMSM is a random copolymer of isobutylene and p-methylstyrene, with reactive benzyl bromide functionality provided by free-radical bromination. The reactive benzyl bromide functionality, C6H5CH2Br is introduced by the selective free-radical bromination of the methyl group of the pendant methylstyryl group on the copolymer. This functionalized terpolymer will maintain polyisobutylene properties such as low

TEMPERATURE AND COMPOUND PERMEABILITY

355

TABLE 11.6 Air Oven Aging of Bromobutyl Rubber, Exxpro, and Nanocomposite Model Tire Innerliner Compounds [6] Compound

3

BIIR 2222 Nanocomposites base polymer (BIMSM) Nanocomposite (with nanoclay) N 660 Naphthenic oil Struktol 40MS Phenolic Tackifying Resin Stearic acid Zinc oxide MBTS Sulfur

4

5

100.00 100.00 60.00 8.00 7.00 4.00 1.00 1.00 1.25 0.50

60.00 8.00 7.00 4.00 1.00 1.00 1.25 0.50

100.00 60.00 3.50 7.00 4.00 1.00 1.00 1.25 0.50

53.3

52.0

64.6

Cure rheometer (160 C) DTorque (MH-ML) (dN m) t10 (min) t90 (min)

3.4 2.1 10.6

5.5 2.9 7.7

5.9 3.6 12.1

Vulcanization rate (dN m/min) Tensile strength (MPa) Elongation (%) 300% modulus (MPa) Hardness A (Shore A) Tear strength (KN/m)

0.8 7.7 795.0 2.6 43.0 35.7

1.5 7.5 815.0 4.0 51.0 41.3

0.9 7.6 755.0 4.9 56.0 45.3

Permeability

Air Oven Aging

Unaged original samples Test temperature (
C) Permeation coefficient (cc mm/(m2-day)) Permeability coefficient (cc mm/(m2-day-mmHg))

40 219 0.322

40 184 0.271

40 133 0.195

Aged 7 days at 100
C Test temperature (
C) Permeation coefficient (cc mm/m2-day) Permeability coefficient (cc mm/(m2-day-mmHg))

40 209 0.308

40 176 0.258

40 123 0.181

Aged days at 125
C Test temperature (
C) Permeation coefficient (cc mm/(m2-day)) Permeability coefficient (cc mm/(m2-day-mmHg))

40 167 0.245

40 140 0.205

40 113 0.166

Mooney viscosity ML1 þ 4 (MU)


356

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

Rheometer (MDR) at 180ºC 7.20

Nanocomposite

6.40

BIMSM

5.60 4.80 4.00

Bromobutyl

3.20 2.40 1.60 0.80 0.00

3.33

6.67

10.00

13.33

15.67 Time

FIGURE 11.7 Vulcanization (MDR rheometer) profiles for bromobutyl, BIMSM, and nanocomposite innerliner formulations [6].

permeability and good vibration damping, while increasing the resistance to oxidative, heat aging reversion resistance. The stability of BIMSM and nanocomposite formulations can be demonstrated when viewing the MDR rheometer profiles at 180
C (Figure 11.7). While bromobutyl compounds will eventually show reversion at extended times at high temperature, the BIMSM formulations will show essentially a level or plateau phase with no reversion. 11.5 VULCANIZATION OF NANOCOMPOSITE COMPOUNDS AND PERMEABILITY Usuki et al. [17] studied the effect of various accelerators in the cure system of an EPDM octadecylammonium-treated montmorillonite clay compound. The specific accelerators investigated included ethylthiourea (ETU), mercaptobenzothiazole (MBT), cyclohexylbenzothiazole sulfenamide (CBS), tetramethylthiuram disulfide (TMTD), and zinc dimethyl dithiocarbamate (ZDMC). Other materials used in the formulation included zinc oxide and stearic acid. The authors reported the following nitrogen permeability results (Table 11.7). Compounds with the lowest permeability (B and C) contained ZDMC and TMTD. Compound F with ethyl thiourea also had a low permeability. The authors attributed the performance of the ZDMC and TMTD compounds to the introduction of pendent accelerator residues along the polymer chain. This introduced polarity to the nonpolar polymer, thereby facilitating an improvement in the clay

VULCANIZATION OF NANOCOMPOSITE COMPOUNDS AND PERMEABILITY

357

TABLE 11.7 Influence of Cure System Accelerator on N2 Permeability [17] Compound

Clay Content Accelerator N2 permeability (10  9 cm3 cm/(cm2 s cmHg))

A

B

C

D

E

F

No ZDMC 2.4

Yes ZDMC 1.7

Yes TMTD 1.7

Yes CBS 2.1

Yes MBT 2.2

Yes ETU 1.9

intercalation or exfoliation state. This work had been done in gum stock containing 100 phr EPDM, 5 phr ZnO, 1.0 phr stearic acid, 1.5 phr sulfur, and 1.5 phr of accelerator. It was noted that the ETU compound also displayed low permeability. ETU would not readily form pendent groups such as what is found with MBT or thiuram derivatives [18]. However the scorch resistance of thiourea accelerators is much shorter than thiuram or dithiocarbamate accelerators that would infer other factors might also influence the morphology of a nanocomposite. Thiourea accelerators also may require the presence of a primary accelerator such as MBTS. Therefore to elaborate on this work a series of accelerators where evaluated in the model innerliner compound in Table 11.8 but where isobutylene-co-para-methylstyrene (BIMSM) and an organoclay was used in place of bromobutyl [8]. Compounds were prepared by melt mixing in a laboratory internal mixer and the accelerators included MBTS, CBS, ZDMC, TMTD, DTDM (4,4-dithiodimorpholine), MBT, DBTU (dibutylthiourea), and the phosphate, ZDBP (zinc dibutylphosphorodithiate). DBTU was selected in place of ETU because it is reported to have superior reversion resistance performance [18]. The lowest permeation coefficients were achieved with TMTD and DBTU that is in agreement with the data presented by Usuki and coworkers (Figure 11.7). The highest permeation coefficients were found with ZDPD possibly due to the lower polarity of any pendent groups, if any, that might be formed along the isobutylene p-methylstyrene copolymer chain [8] (Figure 11.8). Thioureas are known to provide good reversion resistance. The rheometer profile for the nanocomposite compound, bromobutyl control compound, and nanocomposite compound with DBTU in the cure system are shown in Figure 11.9. DBTU and TMTD have the shortest scorch or induction times while the accelerators with the longest induction times representing compounds with the highest permeability. Elaborating on the selection of DBTU, a comparative study had been conducted earlier evaluating a range of thiourea accelerators [18]. Of those studied, DBTU cure systems were found to show little reversion. Compared to the aromatic derivatives, DPTU (diphenylthiourea) and DOTTU (diorthotolylthiourea), DBTU also demonstrated a higher state of cure (DT). DBTU was therefore selected for use in the nanocomposite compound study.

358

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

TABLE 11.8 Properties of Solution Prepared Nanocomposite Compound Compound BIIR 2222 Nanocomposite Carbon black N660 Naphthenic oil Struktol 40MS SP-1068 Stearic acid ZnO MBTS Sulfur DBTU Mooney viscosity ML1 þ 4 (MU) Mooney scorch, t5 (min) MPT-capillary rheometer Shear rate (s  1) Shear stress (kPa) Viscosity (kPa s) Run die swell (%) Cure rheometer (160
C) DTorque (MH-ML) (dN m) t10 (min) t90 (min) Vulcanization rate (dN m/min) Tensile strength (MPa) Elongation (%) 300% modulus (MPa) Hardness A (Shore A) Tear strength (ASTM Die B) (kN/m) Permeation coefficient (cc mm/(m2-day))

6

7

8

60.00 8.00 7.00 4.00 1.00 1.00 1.25 0.50

100.00 60.00 3.50 7.00 4.00 1.00 1.00 1.25 0.50

100.00 60.00 3.50 7.00 4.00 1.00 1.00 1.25 0.50 1.00

55.4 24.1

61.1 19.6

85.1 6.6

101.7 256.0 2.5 12.4

101.7 275.8 2.7 6.1

101.7 819.6 8.1 28.1

3.5 1.7 9.9 0.7

4.6 2.9 20.6 0.3

3.9 1.1 12.4 1.2

9.8 801.3 3.1 47.2 49.3 198.3

9.9 754.2 5.0 62.7 52.3 136.3

10.5 678.1 4.8 58.1 53.1 120.3

100.00

11.6 THERMODYNAMICS AND BIMSM MONTMORILLONITE NANOCOMPOSITES A nanocomposite can be prepared in solution consisting of BIMSM and nanoclay. A set of compounded properties are tabulated in Table 11.8. The nanocomposite compound when compared to the bromobutyl compound showed the following: 1. An Increase in Mooney Viscosity. Using a capillary rheometer shear stress and viscosity results similarly showed a significant increase in viscosity. 2. Scorch or induction time was reduced though cure time was satisfactory.

THERMODYNAMICS AND BIMSM MONTMORILLONITE NANOCOMPOSITES

Permeation coefficient rating (BIIR =100)

62

359

61 60

60 58

58

58

58

58

57

57

56

55

54 53 52

52 50 48 46 MBTS

MBT

CSS

DBTU TMTD ZDEC

ZDBC TDEC DTDU

ZEPD

HTM

Note: ZDEC, zinc diththyl dithiocarbamte; ZBEC, zinc dibenyldithiocarbamate; TDEC, telhrium diethyldithiocarbamater; HTM, hexamethyltetramine.

FIGURE 11.8 Permeation coefficient rating and effect of primary accelerator [8].

3. Tensile strength, modulus, elongation, tear strength, hardness were satisfactory though in the case of Compound 8 the elongation started to drop below 700% [8]. 4. Permeability was reduced by 40%. Viewing the rheometer profile and the data on Table 11.8, the nanocomposite compound containing DBTU (Compound 8) clearly has a shorter induction time

7.20 6.40

Nanocomposite (Compound 7)

Nanocomposite with DBTU (Compound 8)

5.60 4.60

Bromobutyl (Compound 6)

4.00 3.20 2.40 1.60 0.00 0.00

10.00

30.00

20.00


40.00

50.00 Time

FIGURE 11.9 Rheometer profile at 160 C of solution prepared nanocomposite with and without DBTU [8].

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PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

compared to Compound 7 or the bromobutyl rubber compound control [18]. This could infer a relationship to permeability performance. A similar comment was made by Liang et al. [16]. Using the Gibbs free energy equation, the energy change associated with the separation of clay layer plates dispersed in a polymer, DG, can be interpreted to consist of two elements, that is, an internal energy change associated with the formation of new molecular interactions, DE, and an entropy change associated with the conformation of the nanocomposite constituents, DS. The authors therefore used the Gibbs free energy equation: DG ¼ DE  T DS

ð11:9Þ

where T is temperature in K, to describe the energy change accompanying a change in clay silicate gallery spacing from an initial value described as ho to a new distance hf. If DG < 0, it suggest that the silicate gallery height change is exothermic and is favorable. Should DG > 0, then the change in clay plate dispersion is endothermic or not favorable. However, the reverse process would be favored. When DG ¼ 0, then the system will be at an equilibrium with a random distribution of gallery spacing between hf and ho. The entropy change, DS, is a summation of the entropy change of the surfactant molecules (DSsurfactant) and the entropy change of the polymer molecules (DSpolymer) caused by the change in gallery distance, hf  ho, that is DS  DSpolymer þ DSsurfactant

ð11:10Þ

Compared to many plastics and also other elastomers, butyl rubber has a high molecular weight. Viscosity and shear stresses for rubber compounds during mixing are very high and molecular chains can be orientated by such shear stress both in an internal mixer and particularly in a two roll mill operation with a consequent decrease in the entropy of the polymer. Such shearing of stacked nanoclay silicate plates during mixing will separate these plates with polymer chains being inserted between them (i.e., the silicate gallery spacing). Liang et al. reported that this type of structure is thermodynamically unstable [16]. However if such intercalated structures are formed through mixing of butyl rubber and a nanoclay, the transformation to a more stable structure would be prevented due to the very high viscosity of the nanocomposite at room temperature or rapid initiation of cross-link formation, again with increase in viscosity. When the temperature increases with a corresponding rapid decrease in viscosity, then the mobility of the rubber chains can increase and the energy barrier to structural transformation would be reduced. A new intercalated structure with a lower gallery spacing and reagglomeration of the clay can then occur. Extended periods at elevated temperature in a low viscosity, unvulcanized state will facilitate clay plate reagglomeration to the detriment of properties such as permeability [8,16]. Therefore, if such formulated compounds have short vulcanization induction periods and faster cure rates, then the aggregation phenomenon might be arrested and a wide nanoclay gallery spacing and greater nanoclay plate distribution through the

THERMODYNAMICS AND BIMSM MONTMORILLONITE NANOCOMPOSITES

361

FIGURE 11.10 Nanocomposite TEM micrographs illustrating the development of exfoliated clay plates [15].

polymer matrix may be maintained. Elaborating on earlier discussions, Figures 11.10 illustrate this for the three compounds described in Table 11.8. Figure 11.10 shows TEM micrographs for the nanocomposite gum nanocomposite. At 20 nm the clay plates are clearly visible. The micrographs show that the nanoclay is very well dispersed and an intercalated-flocculated condition may have been achieved. Figure 11.10 illustrates the nanoclay dispersion in a fully formulated compound (Table 11.8, Compound 7).

362

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

Montmorillonite clay plates have a nominal diameter in the order of 70–120 nm. The carbon black, grade N660, shown in these micrographs and which is typically used in innerliner compounds has a particle size of 50–80 nm. Aggregates of particles will have much larger dimensions. Therefore, should a clay plate stack become associated with a carbon black particle or aggregate, the contribution to increase in the tortuous path that an oxygen or nitrogen molecule will follow may be negated. Carbon black can have polar functional groups on the surface. Such functional groups can include hydroxyl groups, aldehydes, quinines, ketone, and carboxylic acids. These functional groups in carbon black will influence compound properties such as cure rate and may even influence parameters such as embedded wire corrosion in tires. Rheometer vulcanization testing for effect of carbon black on compound cure rate has been used as a quality assurance test, and is required in many tire company raw material purchase specifications for this reason (i.e., impact on tire compound cure rates). The micrographs above also show that silicate plates can become associated with the carbon black surface, possibly due to the polarity on the silicate surfaces. When DBTU is added to the compound vulcanization system the compound scorch time has been reduced from 19.6 to 6.6 min (Mooney scorch t5 at 125
C). The mooney viscosity has also increased. If the above interpretation is correct, shorter scorch times and increased viscosity will hinder the reagglomeration of silicate plates or aggregates. Figure 11.10 shows TEM micrographs of Compound 8 in Table 11.8. It can be seen that there is an apparent decrease in the amount of clay associated with carbon black particles. At the onset of cure (induction period) the temperature of the compound is increasing and the viscosity decreases as the mobility of chains increase. The energy barrier to structural rearrangement of clay plates can decrease and newly intercalated structures with smaller gallery heights can form. However with short cure induction periods, viscosity reduction can be minimized, thereby maintaining the separation of the clay plates, maintenance of larger gallery spacing, and prevention of carbon black–clay plates associations. More clay is thus available to participate in the formation of a “tortuous path” as illustrated in Figure 11.2.

11.7 NANOCOMPOSITES AND TIRE PERFORMANCE To demonstrate the practical application of nanocomposite in tires, heavy-duty truck tires were built with both conventional and nanocomposite innerliners. Because of the improved permeability of the nanocomposite innerliner compound, the gauge of the innerliner was reduced by 35% compared to the control tire [10]. On completion of the build, tires were visually inspected, a cut tire analysis conducted, and then tested for uniformity that included balance, run out, conicity, and force variation. All tires were X-rayed and inspected for internal separations using shearography. It was reported that tire inflation pressure retention properties were better than the laboratory prepared sample test data would have predicted. Highlights of the reported results included the following:

NANOCOMPOSITES AND TIRE PERFORMANCE

. .

.

.

.

363

The nanocomposite innerliner gauge reduction resulted in an average tire weight savings of over 1.0 kg. Viewing the quality of the nanocomposite tire, whole tire X-rays showed no detriment in parameters such as ply wire spacing, and shearography did not show internal separations in the casing. Liner splice integrity and adhesion for both the bromobutyl innerliner and nanocomposite liner were equivalent. Green tires were also aged in the factory before curing. No liner or liner splice openings were found. To maintain equal output rates at the innerliner extrusion line, extruder RPM speed for the nanocomposite innerliner was reduced 40% with resulting energy savings. The reduced nanocomposite innerliner gauge allowed a tire cure time reduction of 2%.

Tires were tested for uniformity, inflation pressure retention, and durability. Tire uniformity such as balance, run out, force variation, and conicity for the two sets of tires were equivalent. Even though there is a reduction in the liner gauge, at 94 days the average inflation pressure retention of the nanocomposite tire, expressed as a percent loss, is superior to the bromobutyl innerliner tire. This correlated directly with tire durability. After 85 days, the intracarcass pressure for the bromobutyl tire was 113 kPa compared to 108 kPa for the nanocomposite innerliner. Durability was assessed by conducting a simple dynamometer test to FMVSS 119. Rather than stopping at 47 h, the test continued running at 50 kph and the load on the tire was increased by 10% at 8 h increments until the tires failed. Tire temperature was measured at the tread centerline and shoulder at each load step increase. The following observations were noted: . .

.

At 47 h the running temperature of the nanocomposite tire, measured by a needle probe, was 13
C cooler than the tire with the bromobutyl innerliner. Tire test duration for the bromobutyl control and nanocomposite tires, in ratings relative to the control at 100, was 100 and 125, respectively. The ratings for mileage to failure were 100 and 123, respectively (higher ratings are better). The tire load at failure, in ratings, was 100 and 116 for the bromobutyl and nanocomposite tires, respectively.

The potential benefits in truck tire retreadability by use of nanocomposites could be extrapolated from this initial durability data and thus would merit further study. Furthermore, the benefits of nanocomposite innerliners may be more extensive than this preliminary data suggests. For example, the effect on tire rolling resistance, field durability that would be assessed from market trials, tread wear, and overall inservice miles-to-removal merit investigation. Improvement in innerliner performance, in the order of magnitude demonstrated here presents a number of opportunities in further improving tire performance. As noted earlier, reduction in permeability leads to improved tire inflation pressure

364

PERMEABILITY OF RUBBER COMPOSITIONS CONTAINING CLAY

retention. In many instances, this parallels improvement in ICP that correlates favorable with improved tire dynamometer durability [9,10]. Alternatively, lower permeability may allow a corresponding reduction in tire innerliner gauge and a reduction in tire weight. Reduced gauges will favorably influence tire operating temperatures. Reduced innerliner gauge may also afford improvements in factory mixing equipment utilization and efficiency, offer potential increase in liner compound calendering or extruder output rates, due to thinner gauges, and potential cure time savings. 11.8 SUMMARY Nanocomposites offer the opportunity to further improve the performance of tire compounds not possible by other methods. In the case of tire innerliners, bromobutyl rubber compounds have shown to have the best properties with regard to permeability, aging resistance, and fatigue and crack growth resistance. Nanocomposites however can allow a further improvement in innerliner performance. Innerliner elastomer nanocomposites essentially consist of the base polymer and a nanoclay that is typically based on an organically montmorillonite clay. The optimum organoclay content is in the order of 10 phr, beyond which there may be a deterioration in compound mechanical properties. Furthermore, solution prepared nanocomposites are much more effective than melt mixed compositions. BIMSM is much more efficient in achieving a nanocomposite than bromobutyl rubber. REFERENCES 1. Waddell, W. H.; Tsou, A. H. butyl rubber. In Rubber Compounding, Chemistry and Applications, M. B. Rodgers (ed.), Marcel Dekker, Inc., New York, 2004. 2. Tse, M. F.; McElrath, K. O.; Wang, H. C. Relating De Mattia cut growth to network structure of crosslinked elastomers. Polym. Eng. Sci., 42 (6), 1210–1219 (2002). 3. Solis, S.; Rodgers, B.; Tambe, N.; Sharma, B. B.; Waddell, W. H. A review of the vulcanization of isobutylene-based elastomers. Presented at a Meeting of the American Chemical Society Rubber Division, San Antonio, TX, 2005. 4. Van Amerongen, G. J. The effect of fillers on the permeability of rubber to gases. Rubber Chem. Technol. 28 821–832 (1955). 5. Van Amerongen, G. J. Diffusion in elastomers. Rubber Chem. Technol. 37, 1065–1152 (1964) 6. Rodgers, B.; Jacob, S.; Sharma, B. B.; Manjunath, H.; Pal, S. Effect of aging on the permeability of bromobutyl based tire innerliner compositions. Paper 52. Presented at a Meeting of the American Chemical Society, Pittsburgh, 2009. 7. ASTM, D3985. Standard Test Method for Oxygen Transmission Rate through Plastic Film and Sheeting Using a Coulometric Sensor, 2005. 8. Soisson, J.; Rodgers, B.; Weng, W.; Webb, R.; Jacob, S. Vulcanization of nanocomposite tire innerliner compounds and permeability. Paper 112. Presented at a Meeting of the American Chemical Society, Pittsburgh, 2009.

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9. www.butylrubber.com, 2010. 10. Rodgers, B.; Webb, R.; Weng, W. Advances in tire innerliner technology. Presented at a Meeting of the American Chemical Society Rubber Division, Pittsburgh, 2005. 11. Ray, S. S.; Okamoto, M. Polymer/layered silicate nanocomposites: a review from preparation to processing. Prog. Polym. Sci., 28, 1539–1641 (2003). 12. Varghese, S.; Karger-Kocsis, J.; Panikottu, A. Rubber nanocomposites via solution and melt intercalation. Rubber World, 230, 32–38 (2004). 13. Neilsen, L. E. Models for the permeability of filled polymer systems. J. Macromol. Sci. (Chem.), A1 (5) 929–942 (1967). 14. Bharadwaj, R. K. Modeling the barrier properties of polymer-layered silicate nanocomposites. Macromolecules, 34, 9189–9192 (2001). 15. ExxonMobil data. 16. Liang, Y. R.; Ma, J.; Lu, Y. L.; Wu, Y. P.; Zhang, L. Q. Effects of heat and pressure on intercalation structures of isobutylene–isoprene rubber/clay nanocomposites, 1, prepared by melt blending. J. Polym. Sci. Part B. Polym. Phys., 43, 2653–2664 (2005). 17. Usuki, A.; Tukigase, A.; Kato, M. Preparation and properties of EPDM–clay hybrids. Polymer, 43, 2185–2189 (2002). 18. Rodgers, M. B.; Solis, S. C.; Tambe, N. K.; Sharma, B. B.; Waddell, W. H. Thiourea accelerators in the vulcanization of butyl elastomers. Presented at a Meeting of the American Chemical Society, Rubber Division, San Antonio, 2005.

SECTION III

COMPOUNDS WITH RUBBER–CLAY NANOCOMPOSITES

CHAPTER 12

RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER KONSTANTINOS G. GATOS JÓZSEF KARGER-KOCSIS

12.1 INTRODUCTION Many of the rubber applications are based on diene rubber compounds. Their wide use is mostly due to their versatile and highly unsaturated macromolecular backbone providing operational freedom in blending with other rubbers and selecting the desirable vulcanization system. Going back to nineteenth century, the first rubber reported in industrial applications was cis-1,4-polyisoprene widely known as natural rubber (NR), which is a institutive member of this category. Its synthetic substitute, named isoprene rubber (IR, polyisoprene) comforted the crescent raw material prices and availability. The polymerization of butadiene at the beginning of the twentieth century delivered a very popular diene-rubber grade, viz. polybutadiene rubber (BR). This type is widely used in tire industry, mainly due to its low rolling resistance [1]. Apart from the homopolymers described above, the copolymerization of wellselected monomers resulted in important engineering rubbers. By copolymerizing butadiene and styrene monomers through emulsion or solution polymerizations styrene–butadiene rubbers (SBR) were obtained. Its convenient production, versatile structure (e.g., styrene–butadiene ratio) and relative low price contributed to the widespread applications of SBR in different fields. The copolymerization of isobutylene with isoprene monomers was achieved at the beginning of 1940s, and the product marketed under the name of butyl rubber (IIR). This rubber type became popular mainly due to its low gas permeability. As IIR itself is not reactive enough for covulcanization with diene rubbers, substitution of one hydrogen atom of the isoprene monomer by a halogen (bromine and chlorine) is required [1].

Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti. Ó 2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

This chapter deals with “apolar” diene rubbers as a separate chapter is dedicated in this book to the polar variants. On the other hand, epoxidized natural rubber (ENR) and maleated rubbers are also treated in this chapter though they belong to the polar diene rubbers. This grouping is reasoned by the fact that both epoxidation and maleation are so called polymer analogous reactions, which—unlike copolymerization—do not affect much the characteristics of the main chain. Traditionally, the mechanical performance of diene rubbers was enhanced by incorporation of active fillers like carbon black and silica. The use of nonactive reinforcements, such as kaolin, dolomite, chalk, served mainly to increase the hardness and to lower the price of the compound [2]. When an effective reinforcement was required by such minerals, they were surface modified with organosilanes. Using aminosilanes, the hydrolyzed silanol group of the modifier reacted with the filler surface while the pendent amino group, at the other end of the molecule, was involved in the vulcanization of the rubber [3]. In the last decade, specific clay minerals, which belong to the category of layered silicates or phyllosilicates, have attracted great academic and industrial interests for diene rubbers’ reinforcement. Their crystalline structure comprise silica layers and alumina sheets joined together in varying proportions and stacked on top of each other in a certain way [3]. Said layers and sheets, which are superimposed in different fashions, incorporate two-dimensional arrays of silicon–oxygen tetrahedral and twodimensional arrays of aluminum– or magnesium–oxygen/hydroxyl octahedral units. Condensation in a 2 : 1 ratio of two silica sheets with an alumina, one gives rise to a 2 : 1 type phyllosilicate (e.g., pyrophyllite, smectite or montmorillonite, vermiculate, illite). Among the 2 : 1 layered clays, the montmorillonite (Mt) is most distinguished because of its significant “swelling” capabilities. This behavior emanates from partial replacement of trivalent Al by divalent Mg in the octahedral sheet resulting in a negative surface charge of the layer. By this way the space between the surface of the layers (i.e., interlayer space) is capable of accommodating cations. When Mt is contacted with water or with water vapor, the water molecules penetrate between the layers (i.e., interlayer swelling), forming a hydration shell around the cations. This increases the basal spacing of the Mt. The charge compensating cations in the galleries may be easily exchanged by other cations when available in aqueous solution; hence they are called “exchangeable.” The total amount of these cations is expressed in milliequivalents per 100 g of dry clay, and called as the “cation exchange capacity” (CEC) of the clay [4]. The reinforcement of diene rubbers by layered clays depends on the dispersion degree of the silicate layers in the rubber matrix. More specific, it is desirable that individual silicate sheets are surrounded by the macromolecular chains of the matrix. In this way the filler–matrix interfacial interactions are fully developed. As shown in Figure 12.1, when the macromolecular chains have been penetrated into the interlayer space of the clay layers, expanding at the same time the basal spacing (d(001)), the so-called “intercalated” structure is obtained. If the clay layers are fully separated apart then one reports on “exfoliated” structure. Apart from these two major categories, several slightly differentiated structures have been proposed in the literature [5]. Considering the fact that the thickness of the said silicate sheets

PREPARATION METHODS

371

FIGURE 12.1 Scheme of the different layered clay dispersions in a rubber matrix.

is in the range of 1 nm, such composites are referred to as diene rubber–layered clay nanocomposites.

12.2 PREPARATION METHODS The clay particles in a diene rubber matrix exist in different dispersion structures. Their characteristics strongly depend on the preparation methods and their conditions. 12.2.1 Latex The hydrophilic character of clay renders it suitable for its mixing with diene rubber latices. Latex is considered as an aqueous dispersion of fine rubber particles in the submicron–micron range. Generally speaking, swellable phyllosilicates are capable of intercalating water through which the interlayer space increases to about 2 nm, due to the hydration of the interlayer cations. At this stage a further expansion is not favored. In pristine clays with small monovalent cations, the basal spacing increases up to 4 nm, or even further, in presence of excess water. Frictional forces rising from edge-to-face particle association leads to paste-like clay/water dispersions. On stirring, the individual silicate layers get separated by large distances, which is limited only by the volume of water in the dispersion [3]. In order to assist further the dispersion of clay in water, sonication methods can be applied [6]. Nowadays, a plethora of rubbers are available in latex form [7]. Therefore, mixing aqueous clay slurry with rubber latex is considered as a convenient and cost-effective process for rubber reinforcement. As NR is by far the most traditional rubber latex,

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

extended research has been conducted on NR/clay systems. Aqueous dispersion of sodium fluorohectorite (FHT, a synthetic layered silicate) with 10 wt.% solid content was first prepared in a mechanical stirrer followed by its mixing with NR latex. Dispersions of sulfur type curing ingredients were added to the mixture and left to mature. By casting the compound in the desired mould and after its vulcanization, the FHT layers were finely dispersed in the NR matrix [8]. In order to accelerate the water evaporation, the rubber latex/clay dispersion was concentrated in an evaporator under reduced pressure. This method was adapted for NR/Mt suspension wherein the Mt layers were finely dispersed even at an amount of 30 parts per hundred rubber (phr) [9]. To accelerate curing, NR may be prevulcanized by sulfur and the related latex to be stabilized by ammonia. Aqueous solution of sodium FHT was stirred with the prevulcanized NR latex and the compound was cast and left to fully vulcanize in an oven. The related process led to intercalated instead of exfoliated structures [10]. This structure was favored by prevulcanized NR particles having restricted chain mobility. Note that the NR particles consists of about 86% cis-1,4-polyisoprene, 10% water, 3% lipids, 1% proteins plus several other organic components in very low amounts [1]. In order to assist the coagulation of the rubber/clay suspensions, flocculants are used. For NR/Mt and SBR/Mt an electrolyte solution of 2 wt.% dilute triethylenetetrammonium chloride was used as flocculant [11]. Addition of flocculants causes usually a reaggregation of the exfoliated silicate layers. Exfoliation could be triggered by using latices with very small rubber particles [11]. Figure 12.2 gives a scheme of the preparation of rubber–clay nanocomposites via the latex coagulation route.

After drying under reduced pressure

Latex compound after high stirring

Adding flocculant

After drying under reduced pressure

Latex particle Clay layer Flocculant

FIGURE 12.2 Schematic representation of rubber–clay nanocomposites prepared by the latex method with and without flocculant addition.

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373

For industrial applications, it is often required to blend expensive rubber latices with low-cost versions. This has been presented for polyurethane latex (PUR), which was blended with prevulcanized NR latex followed by their mixing with an aqueous suspension of sodium fluorohectorite [10]. It was found that the silicate layers were more compatible with the polar and unvulcanized PUR phase rather than with the NR in prevulcanized stage [12]. However, in blends, preparation of exfoliated clay structures has been also reported. In such an example, NR and SBR latices were mixed at a 50/50 weight ratio and stirred for 30 min followed by the addition of an Mt aqueous slurry (Mt content 2 wt.%). After 30 min of stirring the mixture was coagulated by dilute solution of sulfuric acid (2 wt.%), washed with water and dried at 80
C for 24 h under reduced pressure [13]. 12.2.2 Solution All commercially available diene rubbers are soluble in suitable solvents. Therefore, solution mixing represents an effective way to prepare diene rubber–clay nanocomposites. In order to render the hydrophilic clay hydrophobic and thus compatible with organic solvents, their surface treatment is necessary. Usually an amine salt or a quaternary ammonium salt is added to the clay/water suspension so that the corresponding organic cations replace the previous “occupants” (i.e., inorganic cations) of the clay surfaces. By this cation exchange, the amino group of the modifier (intercalant) is tethered on the clay surface while its hydrocarbon tail is positioned in the interlayer space whereby excluding the water molecules. The result is an organo-modified clay that is called as “organoclay.” This organophilic modification (cation exchange) is associated at the same time with an increase in the basal spacing of the layered clay (Figure 12.3). Toluene has been proved to be a suitable solvent for preparing NR–organoclay nanocomposites [14,15]. NR along with the related sulfur curatives were dissolved in toluene followed by the addition of Mt modified by rather nonpolar dimethyldihydrogenated tallow ammonium or dimethyl hydrogenated tallow (2-ethylhexyl) ammonium salts. After vigorous stirring of the compound, the solvent was evaporated and the sample was dried under vacuum prior vulcanization, which resulted in

Clay

Organo-modification

Organoclay Silicate layer

– NR3+

–

NR3+ d (001)

– + Na –

Silicate layer

– + Na Silicate layer – +

Na

+

Na

d (001)

NR3+ –

Silicate layer

NR3+ –

FIGURE 12.3 Effect of the organophilic modification of clay schematically.

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

intercalated structures [16]. To support the NR swelling/dissolution in toluene the NR was previously masticated whereby its mean molecular mass was reduced [15]. For BR compounds, facile solvent for mixing Mt modified with dimethyl ditallow ammonium with the rubber was a mixture of n-hexane and cyclohexane [17]. Moreover, toluene served as a suitable medium for mixing BR with a dispersion of octadecyl amine modified Mt in ethyl alcohol. The drying of this solution took place in a vacuum oven at 50
C for 2 days [18]. Using the same technique and adding dicumyl peroxide (DCP) additionally to produce BR/organoclay vulcanizates, a similar morphology as found for the noncured version, was reported [19]. Epoxidation of NR, known as ENR, renders the macromolecular chain polar. Mt modified with octadecyl amine was stirred with milled ENR in toluene. The solution was left to dry at room temperature producing intercalated structures [20]. Intercalated structure was also detected when ENR was dissolved in methyl ethyl ketone (MEK) followed by incorporation of very high amount (i.e., 100 phr) of Mt modified with a quaternary ammonium intercalant [21]. In case of SBR, the rubber was first dissolved in toluene wherein DCP was introduced later on. The used Mts were modified with amines bearing various alkyl tail lengths. The latter organophilic Mts were dispersed in ethyl alcohol and each of them was further blended with the SBR rubber solution. The samples, which were obtained after evaporating the solvents in an oven, presented intercalated morphology [22]. Toluene has been used in solution-mixed IIR–organoclay nanocomposites, too. The organo-Mt was dispersed in toluene by vigorous stirring, wherein IIR/toluene solution was added, followed by a 24 h high-speed stirring. The solvent was removed in an oven at 40
C for 24 h and the specimens, as produced, possessed intercalated structures [23]. Blending of rubbers in the appropriate solvent medium followed by incorporation of organoclay solution has been presented for the system of NR/SBR/Mt modified with alkylammonium salt. In that case the common solvent was toluene [24]. The solution mixing technique has been proved feasible to prepare diene rubber–organoclay nanocomposites, either with (ready for vulcanization) or without curatives (introduction of the vulcanization agents occurs usually by melt blending). 12.2.3 Melt Blending A convenient and widely practiced way to produce diene rubber–clay nanocomposites is melt blending. This involves the mixing of the related components in traditional rubber compounding equipments, such as inner mixer and open mill. The availability of such compounding facilities was the driving force to explore the melt blending as production method for rubber–clay nanocomposites. For NR/organoclay vulcanizates usually all mixing components are added on a roll mill at room temperature [25]. In an example, NR was mixed with vulcanization additives on an open two-roll mill, followed by the incorporation of Mt modified with octadecylamine (Mt-ODA). Finally, sulfur was added [26,27]. Incorporation of Mt modified with tetraoctyl phosphonium salt delivered intercalated nanocomposite

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375

structures opposed to Mt modified with triphenyl vinylbenzyl phosphonium salt, which resulted in a microcomposite. This underlines the effect of the aromatic modifier, which discourages the NR macromolecular chains to penetrate into the interlayer space due to steric effects during melt blending [28]. A rather exfoliated dispersion of fluorohectorite modified with octadecylammonium salt has been produced by mixing NR along with the accelerator additives for 10 min on a two-roll mill set at 100
C. The sulfur was added on the open mill at room temperature [29]. NR has been mixed with organoclays in internal mixers, too. In the related example, the compounding took place at 120
C for 10 min at a rotor speed of 60 rpm [30]. In another work, a NR/Mt batch was first obtained through the latex route. The compounding continued in an internal mixer for 5.5 min at 50
C and 50 rpm rotor speed. Therein, the processing ingredients were incorporated. Finally, the vulcanization curatives were added on a two-roll mill [31]. In another type of mixing pattern, the inner mixer served as “reactive chamber.” First, the NR was masticated for 1 min at 80
C followed by the incorporation of Na-Mt and the ammonium salt in an amount corresponding to the CEC of the clay. The ammonium salt itself was dispersed at ca. 70 wt.% in a water/propanol (68/32) mixture. Mixing of the compound was carried out for 5 min and the batch was discharged at 115
C [32]. Here the NR served as dispersion medium for the cationic exchange reaction yielding finally intercalated clay structures. To prepare a similar compound, first the pristine Mt was reacted with ammonium salt in the NR “medium” in an internal mixer. Compounding took place for 5 min at 145
C. The curatives were introduced afterwards on an open mill [33]. For the preparation of IR–organoclay nanocomposites a corotating twin-screw extruder (screw diameter: 40 mm, screw length/diameter ratio: 48), operated at 200 rpm was also adapted. The organoclay was incorporated through the first downstream side feeder, wherein the temperature of the barrel was below 200
C [34]. The organoclays with relative high initial basal spacing, which were mixed with the IR matrix, did not present any further interlayer expansion. However, the shear forces induced during extrusion resulted in a good dispersion of silicate stacks in the IR matrix [34]. Similar results were obtained for the in situ organophilic modification of Mt and its melt blending with IR in an inner mixer. In this operation, IR was the reaction medium for the cation exchange reactions replacing Naþ by di (hydrogenated tallow)-dimethylammonium or hydrogenated tallow-benzyl-dimethylammonium cations (intercalants) [32]. Also BR served as medium for the cationic exchange reaction of pristine Na-Mt with dimethyl dihydrogenated tallow ammonium chloride, as intercalant, in an inner mixer. The melt blending conditions were optimized at 80
C for 10 min mixing time with rotor speed at 150 rpm. Sulfuric curatives were incorporated in a follow-up step (at 30
C at 60 rpm rotor speed for another 5 min). The resultant batch was further compounded on a roll mill for 10 min at room temperature (RT) prior to vulcanization. This procedure yielded intercalated clay structures [35]. This morphology outcome was maintained even up to 50 phr clay loading though some agglomerate could not be avoided [36]. For transparent BR–organoclay nanocomposites, produced by melt blending, DCP has been used as curing agent [37].

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

BR–organoclay nanocomposites were usually prepared by melt blending of the BR with the organoclay in an internal mixer at elevated temperatures followed by the addition of the curatives either in an internal mixer [38] or on an open mill [39,40] at RT. These operations usually resulted in intercalated structures. In order to create exfoliated layered clay dispersion in BR additional compatibilizers have to be added. More specific, BR and Mt modified with quaternary trimethylstearyl ammonium salt were placed in an internal mixer along with the preferred compatibilizers, viz. carboxyl terminated butadiene (CTB) oligomer or carboxyl terminated butadiene acrylonitrile (CTBN) copolymer. Note that both CTB and CTBN are liquid rubbers. The mixing was performed for 10 min, in total, at RT with a rotation speed of 60 rpm. A reblending was performed in the same mixer at 150
C for 4 min followed by the incorporation of the sulfuric curing agents at RT on an open mill [41]. In another example under more reactive conditions, maleinated liquid polybutadiene and organoclay bearing hydroxylethyl groups were mixed in 1,2-dimethoxyethane in order to cause reaction of the active sites. Mixing BR with the organoclay masterbatch, prepared as described above, maintained the intercalated and exfoliated clay structures, which were originally present in the masterbatch [42]. Among the diene rubbers, ENR offers inherently adequate polarity to support the nanocomposite formation with organoclay. For example, ENR has been compounded with Mt modified either with octadecylamine or with methyl-tallow-bis-2-hydroxyethyl quaternary ammonium salt, producing intercalated and exfoliated structures. The mixing of all ingredients, including the sulfuric curatives, was performed on an open mill at RT [43]. For the production of SBR–clay nanocomposites various compounding techniques have been adapted. SBR mixing with organoclay has been performed on a tworoll mill at RT [44,45] or at increased temperatures [29,46]. This resulted in intercalated organoclay dispersion. In another trial, the curatives have been incorporated in an internal mixer. Therein, the increased shear forces developed during compounding induced a characteristic peel-off of the silicate layers in the SBR matrix [47,48]. To favor the clay dispersion in the rubber, the melt blending was preceded by the production of a rubber/clay masterbatch making use of the latex compounding [11,49]. An aqueous slurry of Na-Mt along with interfacial agent and SBR latex were coagulated by dilute sulfuric acid. To this mixture sulfuric-type curatives were further compounded on an open mill for 10 min [50]. In a similar compound, bis (triethoxysilyl-propyl)tetrasulfide (Si69) has been added during compounding on the open mill in order to increase the interfacial interactions in cooperation with the premixed 3-aminopropyltrimethoxysilane agent [51]. Exfoliated structures were also produced by mixing a concentrated suspension of Mt, modified by allylamine, with SBR latex. The curatives were added by melt blending [52]. A solution assisted nanocomposite preparation method has been presented for SBR, wherein the rubber was mixed with Na-Mt dispersed together with dimethyl distearyl ammonium salt in toluene. Addition of Si69 in the compound followed by melt blending in an inner mixer and open mill contributed to an intensive mixing of the swollen SBR with the plain SBR matrix [53].

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For IIR–organoclay nanocomposite preparation, usually direct melt blending on an open mill was practiced that resulted in intercalated structures [54–56]. IIR was also the model rubber to investigate effects of the mixing temperature, time, and pressure on the intercalation kinetics of the organoclay. Note that the compound was free from curatives. It was found that the initially intercalated IIR chains, after blending the IIR and organoclay on an open mill, were partially extracted from the interlayers due to annealing under static conditions even within 10 min at temperatures above 100
C. The reason behind this unstable behavior at high temperature and at atmospheric pressure was traced to the thermodynamically favorable movements of the IIR chains outside of the interlayer space, being associated with entropy gain. Examining the pressure effect in the same rubber stock, it was presented that the increase of pressure during heat treatment resulted in extended clay layer flocculation, and thus in a coarser clay dispersion [56]. An exfoliated dispersion of organoclay was obtained in IIR, however, when using its maleic anhydride (MA) grafted version, viz. IIR-g-MA. The compounding of the IIR-g-MA with the clay bearing octadecylamine (ODA) intercalant took place on an open mill and followed by temperature treatment at 110
C under static conditions [57]. On the other hand, the inherent polarity of bromobutyl rubber (BIIR) was found to support clay intercalation [58]. In many cases, blending of apolar diene rubbers with polar versions proved to be convenient to prepare diene rubber–clay nanocomposites. This has been presented for NR, which was compatibilized with ENR at amounts up to 10 phr [59–61]. With further increasing amount of ENR (up to 50 phr) exfoliated organoclay morphologies were obtained [62]. The size of the ENR droplets in the NR matrix was decreased in presence of the organoclay. At the same time, the silicate platelets were located in the NR/ENR interphase, wherein the organoclay influenced the interfacial tension between the rubber components [62]. In another example, the processing agents used in NR/BR blends improved the dispersion degree of the intercalated clay formations [63].

12.3 CURE CHARACTERISTICS It has been presented for several diene rubber/organoclay systems that the compounding additives, and especially the curative, strongly influence the formation of nanocomposite structures. Monitoring the torque versus time in a curemeter for NR, it has been observed that there is a reduction of optimum cure time when tetraethylene pentamine was added. A rather similar behavior was detected when Mt was incorporated in the compound. Nevertheless, when tetraethylene pentamine was first mixed with Mt at 100
C for 4 h (in order to increase the Mt interlayer space) the optimum cure time of NR was significantly decreased [64]. Incorporation of ODA in NR was associated with a prominent decrease in the cure time. However, the decrease was more pronounced when NR was mixed with Mt modified with ODA (Mt-ODA) than ODA alone, although the same amount of ODA

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

was present in the recipes. This observation was accompanied by an enhanced torque plateau value measured by a curemeter. By means of differential scanning calorimetry (DSC) a decrease in the activation energy, required for vulcanization, was found when Mt-ODA served as filler [65]. For the same rubber/filler (i.e., Mt-ODA) combination and similar sulfuric curing agents a decrease in the scorch time was reported. The accelerating effect of Mt-ODA was traced to the formation of a Zn-based complex during curing [66]. Under such conditions, zinc is prone for coordination complexing in which the amine groups of the intercalant and the sulfur participate. It was further speculated that especially the ODA intercalant may detach from the Mt surface and participate in the complex formation. The migration of the amine intercalant toward the rubber matrix accelerates the vulcanization. This phenomenon is accompanied with deintercalation, that is, collapse of the interlayer space [43,66]. Trapping of the Zn-complexes in the interlayer space promoted the onset of intercalated and exfoliated clay dispersions in NR. On the other hand, curing with high-energy radiation should eliminate any curative-related deintercalation process. This was proved by radiation curing of a NR–Mt-ODA nanocomposite containing low amount of organoclay. As expected, the exfoliation morphology remained and did not turn into intercalated one [30]. Addition of organoclay does not yield always accelerated curing. When an aromatic phosphonium salt served as intercalant for Mt, which did not favor the intercalation of NR, the vulcanization characteristics were not affected. That was not the case when an aliphatic phosphonium salt worked as intercalant, which additionally lowered the activation energy of curing, and increased the cross-link density of the related rubber [28]. Incorporation of compatibilizer in the rubber/organoclay compound does not reduce further the cure time, as it has been shown for BR–organoclay nanocomposites. Adding CTB, the initially highly reduced optimum cure time remained rather unaffected, while further addition of CTB slightly increased the related value [41]. On the contrary, the optimum cure rate and scorch time of BR–organoclay nanocomposites could be further reduced using organoclay with increased CEC. The effects of two Mts, both modified with dimethyl dehydrogenated tallow quaternary ammonium salt, and having different CEC values (125 and 95 mequiv./100 g, respectively) have been compared. The clay of higher CEC, accommodating higher amount of intercalant in the interlayers than the clay of lower CEC, promoted the vulcanization [39]. The exothermic reaction of curing of the BR/organoclay was monitored by means of DSC. It has been found that the catalytic action of the organoclay in the curing reaction resulted in a shift of the exothermic cure peak to significant lower temperatures [40]. The shift was more pronounced for the organoclay with higher CEC and hydrophobicity. Experiments on BR–organoclay nanocomposites revealed the role of the sulfur curatives in this diene rubber. By melt blending of BR with Mt-ODA in an inner mixer resulted in clay intercalation. Incorporation of sulfur type curatives in the compound increased further the initial basal spacing [38]. This fact suggests the presence of curatives also in the interlayer space. Model experiments showed that intercalation

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379

took place also when sulfuric curatives were solution mixed with the organoclay in absence of rubber [55]. Effects of mixing time and rotor speed in an internal mixer on the cure characteristics were studied for a BR mixed with Mt, dimethyl dehydrogenated tallow ammonium chloride, and sulfuric curatives [35]. It was found that with increasing mixing time or rotor speed the scorch time and optimum cure time of the compounds were reduced. Increased cross-link density in the vicinity of the organoclay platelets compared to the bulk has been proposed for SBR mixed with Mt modified by octadecyl trimethylamine [44]. By increasing the organoclay content, the heat of the curing reaction, detected by DSC under isothermal conditions, was also increased. This indicated for a higher degree of cross-linking. Investigating the role of curatives for the same rubber, it was found that the presence of stearic acid did not affect the nanocomposite structure in respect to the interlayer space of Mt modified initially with dimethyl dialkyl ammonium salt [58]. The ability of the sulfuric curatives to intercalate with Mt-ODA has been presented for IIR. Vulcanization of the related stocks at 140, 160, or 180
C, respectively, resulted in both intercalated and deintercalated organoclay populations. When the deintercalation was less the vulcanization temperature was lower [55]. The nanocomposite morphologies, corresponding to exfoliated, intercalated, or deintercalated populations, depend on several factors [67]. Cross-linking inside the clay galleries may cause exfoliation, whereas when the primary amine-intercalant migrates into the rubber matrix to participate in the zinc/sulfur complex formation, deintercalation may be the result [68]. Under given conditions, intralayer crosslinking seems to occur faster than extralayer one [69]. Moreover, rubber curatives may migrate toward the organoclay interlayer space [70]. Exfoliation of organoclay platelets in IIR matrix has been reported when using its maleic anhydride grafted version (IIR-g-MA). After compounding IIR-g-MA, Mt-ODA, and sulfuric curatives, highly intercalated clay populations were obtained. The intercalated morphology turned to an exfoliated one when the compound was cured for 80 min at 150
C [57]. As mentioned before, the inherent polarity of ENR is beneficial for organoclay intercalation. At the same time ENR accelerates the curing process when blended with NR [59]. Increasing epoxidation (from 25 to 50 mol%, referred to ENR25 and ENR50, respectively) supported the clay exfoliation. Blending ENR50 with NR and organoclay resulted in exfoliated structures of the latter in the vulcanizates when the ENR50 content was higher than the NR in the corresponding recipes [62]. Accelerated curing was also noticed for other organoclay filled diene rubber blends. The scorch time of NR/SBR compound was decreased by increasing the amount of the organoclay [24].

12.4 CLAY DISPERSION It has been shown above how important is to select the right preparation method for mixing a diene rubber with layered clays. In addition, each clay modification may

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

trigger specific interactions with the rubber matrix and with further recipe components (e.g., curing agents). Moreover, conditions of the vulcanization process also affect the final dispersion of the clay in the corresponding vulcanizate. This is the reason why the layered clay dispersion in each rubber stock should be detected and characterized accordingly. 12.4.1 Detection To identify clay platelets of about 1 nm thickness in rubber matrix require highresolution techniques, for example, those based on electron beam/material interactions [71]. Among the related methods, transmission electron microscopy (TEM) has been established as a standard tool to detect the clay dispersion. Apart from TEM, scanning electron microscopy (SEM) and atomic force microscopy (AFM) are also frequently used. For TEM investigations ultra-thin sections of the related specimen should be cut with a diamond knife in a cryochamber. For AFM studies, smooth surface samples are required that can be produced even by hot pressing between glass plates. As reported before, NR latex mixed with pristine FHT (latex route) yielded finely dispersed FHT platelets in the NR matrix [8]. Moreover, this water-assisted mixing technique preserved the high aspect ratio of the FHT layers as presented in Figure 12.4. By melt mixing of NR with Mt-ODA, individual organoclay layers along with stacks of clay platelets appear in the matrix [66]. The Mt layers, shown in Figure 12.5, possess a markedly lower aspect ratio than the FHT. This is not only due to their different origin (natural versus synthetic) but also due to the more severe mixing

FIGURE 12.4 TEM picture taken from 10 phr fluorohectorite (FHT) filled NR nanocomposites produced via the latex route [8]. Copyright 2003, with permission from Elsevier.

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FIGURE 12.5 TEM picture taken from NR filled with 10 phr of montmorillonite modified by octadecylamine (Mt-ODA) produced through melt mixing.

conditions during melt blending compared to latex compounding. Clay platelet stacks have been found also in NR melt blended with Mt modified with dimethyl dihydrogenated tallow ammonium salt [72]. Similar nanocomposite morphologies were produced when NR served as a medium for mixing the pristine clay with the intercalant simultaneously [32]. By this way, the cationic exchange reaction, that took place during melt blending, resulted in intercalated Mt layer stacks in the NR matrix. BR also resulted in a dispersion of organoclay layer stacks in the rubber matrix, when the related components were melt-blended in an inner mixer [41]. In case of BR, use of CTB compatibilizer resulted in an increased interlayer spacing as detected by TEM [41]. Usually for the preparation of diene rubber–organoclay nanocomposites an amount of about 10 phr organoclay or less is used. However, the polar ENR chains can penetrate well into the organoclay interlayer space, and thus ENR can be mixed with up to 100 phr organoclay. This was demonstrated by the production of ENR/ organoclay masterbatch through the solution method [21]. As shown in Figure 12.6, the TEM image of the related compound presented a rather uniform distribution of the organoclay in the ENR despite the rubber/clay ratio of 1:1. In many cases, organoclay has been combined with conventional fillers, like carbon black or silica. On the example of SBR mixed with Si69 and synthetic FHT modified with protonated a,v-bisamino-terminated polybutadiene liquid rubber, it was shown that the FHT platelets were bridging the silica agglomerates (Figure 12.7, [73]).

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

FIGURE 12.6 TEM image of ENR containing 100 phr organoclay [21]. Copyright 2009, with permission from Elsevier.

When binary rubber blends are modified with clay, TEM investigation is useful to get information for the miscibility of the rubber components. This has been shown for an NR/PUR blend. Prevulcanized NR latex was mixed with PUR latex in 1 : 1 ratio and filled with 10 phr FHT in situ [10]. As presented in Figure 12.8, the layered

FIGURE 12.7 TEM micrographs of SBR vulcanizates filled with a mixture of 25 phr silica and 5 phr FHT-ATB [73]. Copyright H€uthig GmbH. Reproduced with permission.

CLAY DISPERSION

383

FIGURE 12.8 TEM picture taken from the cast film of a PUR/NR (1/1) blend containing 10 phr sodium FHT produced by latex compounding. Note: sulfur–prevulcanized NR appears dark.

silicate stacks were located at the boundary of the PUR and NR phases. Note that the sulfur prevulcanized NR appears darker than the PUR phase in the related image. Similarly, for NR melt blended with ENR50 and Mt modified by a quaternary ammonium salt, the exfoliated organoclay found to encapsulate the ENR50 particles [62]. Additionally, the reduced size of the ENR droplets signified the role of the organoclay acting as compatibilizing agent. Although many TEM images of high magnification have been presented in the literature, quantitative measurements of the related interlayer spacing are scarce. Nevertheless, the qualitative information deduced by direct methods, like TEM, is essential for the morphology characterization of rubber–clay nanocomposites. 12.4.2 Characterization X-ray diffraction (XRD) scattering is a widely used method to quantify the change in the (organo)clay basal spacing and thus get a deeper understanding on the rubber nanocomposite morphologies (as already discussed more in detail in Chapter 7). The repeated distance of the interlayer space plus the thickness of the clay platelet in a phyllosilicate structure raises a peak at low 2u scattering angles in a XRD spectrum. This structural regularity for unmodified or pristine clays gives a peak in the range of 2u ¼ 6–10
angles, which corresponds through the Bragg equation (cf. Eq. 12.1) to a basal spacing of about 1.5–1.0 nm. nl ¼ 2d sin u

ð12:1Þ

384

RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

where n is the reflection order term, d is the interlayer or basal spacing, l is the wavelength of the beam used, and u is the scattering angle of the beam. After organophilic modification of the clay (which replace the former interlayer cations), the basal spacing increases and the initial peak in the XRD spectrum shifts to lower 2u angles. When clay or organoclay is mixed with diene rubbers, this peak either shifts further to lower 2u angles or disappears. The first indicates intercalation, while the latter suggests exfoliation of the phyllosilicate platelets in the corresponding rubber matrix. For example, XRD spectra taken from vulcanizates after melt mixing of NR with clay and organoclay (in 5 phr) suggested the presence of microcomposite and exfoliated nanocomposite structures, respectively [74]. In other cases, the exfoliated morphology of the clay platelets in the NR matrix was not maintained after mixing with the curatives followed by vulcanization. As presented in Figure 12.9, mixing of NR with the organoclay (Mt modified by dimethyl dehydrogenated tallow quaternary ammonium salt) on an open mill at RT resulted in rather disordered/delaminated structures. This structure was deduced based on the presence of a shoulder at low 2u angles. When sulfur-type curatives were incorporated, a reordering of the platelets took place, which was enforced further by vulcanization [75]. Replacing the Mt with laponite, bearing the same intercalant, and performing the same mixing/vulcanization with NR, it was noticed an interesting effect. Organolaponite, possessing a smaller aspect ratio than Mt, resulted in rather exfoliated clay dispersion compared to the Mt-induced intercalated one. It is noteworthy that XRD spectra, taken from various steps of rubber production, delivered important information on the change in the clay exfoliation/(de) intercalation behavior [68]. Cross-linked

16,000

Intensity (c.p.s.)

Uncross-linked Uncross-linked·

12,000

8,000 Curatives

4,000 Nanoclay

0 0

2

3

4

5

6

7

8

2θ

FIGURE 12.9 XRD patterns for the cross-linked NR–O-Mt nanocomposites (top), uncrosslinked NR–O-Mt nanocomposites (middle), and uncross-linked NR–O-Mt nanocomposites without curatives (bottom) [75]. Copyright 2008, with permission from American Chemical Society.

CLAY DISPERSION

385

Although melt blending of NR with pristine clay results in conventional microcomposites, adapting the latex route for the same components nanocomposites can be produced. When NR latex was stirred with 5 phr Mt and coagulated under reduced pressure, exfoliated nanocomposites were found based on the absence of the expected peak in the related XRD spectrum (Figure 12.10). With increasing clay amount in the mixture, however, intercalated structures formed. This was reflected by the onset of a peak at lower scattering angles than that of the pristine Mt. The initial interlayer spacing of Mt at 1.17 nm was expanded for 1.59, 1.69, and 1.72 nm for 30, 20, and 10 phr clay contents, respectively [9]. Effects of the processing steps on the clay dispersion, assessed by XRD, were investigated on the example of BR/(organo)clay compounds [38]. After mixing of BR with Na-Mt in an inner mixer, followed by incorporation of the sulfuric curatives and vulcanization no significant peak shift could be observed. This indicates that a microcomposite was produced. On the other hand, by mixing Mt modified with the primary amine intercalant at 30 phr in the inner mixer, the initial basal spacing of 2.08 nm was expanded to 2.22 nm. This interlayer distance was further increased to 2.72 nm by incorporating the sulfur curatives. The compounding on the open mill prior compression moulding resulted in a further interlayer expansion for 2.88 nm. After vulcanization, an intercalated nanocomposite structure with a basal spacing of 3.50 nm was obtained. By contrast, the Mt modified with quaternary amine intercalant in 30 phr resulted in a different sequence of interlayer spacing when it was mixed with BR. The basal spacing of the quaternary amine intercalated organoclay increased from 3.50 to 3.87 nm when all curing agents were incorporated. Vulcanization resulted in intercalated nanocomposite

5 phr

Intensity (a.u.)

10 phr

20 phr

30 phr

Pure clay 2

3

4

5

6

7 8 2θ (º)

9

10 11 12 13 14 15

FIGURE 12.10 XRD patterns for the pure clay and NR–clay nanocomposites [9]. Copyright 2005, with permission from Elsevier.

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

structure with a basal spacing of 3.99 nm [38]. The different structural outcomes due to the primary and quaternary amine intercalants of the organoclays are usually attributed to their inherent reactivity [76]. The XRD spectra are generally taken by wide-angle X-ray scattering (WAXS) devices working at low scattering angles (2u from 1–2 to 10
). Thus, information from highly expanded clay galleries is missing. In these cases, small-angle X-ray scattering (SAXS) delivers further useful information on the intercalation. In a related example, SBR was melt blended in an inner mixer with organoclay which was first treated with maleinated liquid polybutadiene to increase its basal spacing [42]. SAXS results showed that the treated organoclay exhibited high basal spacing due to the reaction between the pendent succinic rings of the maleinated polybutadiene and the hydroxyl groups present in the ammonium ions of the organoclay. By mixing 20 phr treated organoclay with the SBR the interlayer spacing expanded to 7.2 nm according to SAXS. Effects of various processing steps on the clay dispersion were investigated by XRD also for diene rubber–clay nanocomposites produced by latex compounding [11]. A 3 wt.% Mt aqueous suspension was stirred with SBR latex, followed by co-coagulation with 2 wt.% dilute triethylenetetrammonium chloride solution and drying in an oven. It was found that the Mt was exfoliated in the aqueous slurry. This morphology was maintained after the addition of the SBR latex. On the other hand, incorporation of the flocculant caused a reordering of the Mt platelets reflected by a peak at an interlayer spacing of 1.35 nm. This situation did not change after drying of the related compound. Moreover, evaporation of the water resulted in an even smaller basal spacing of 1.29 nm. This study made obvious the decisive role of the flocculant in the nanocomposite formation when produced by the latex compounding method [11]. Effects of flocculants on the clay dispersion in SBR–clay nanocomposites were also studied. The composites were produced by latex compounding followed by flocculation with 1 wt.% calcium chloride aqueous solution (Ca2þ ) or 2 wt.% sulfuric acid (Hþ ) solution [77]. The SBR/Mt vulcanizates (note that the sulfuric curing agents were incorporated on an open mill) showed a microcomposite structure presenting a basal spacing of 1.43 and 1.51 nm for Hþ and Ca2þ flocculated systems, respectively. This was attributed to the reaggregation of the Mt platelets induced by the flocculant. In order to overcome this problem, interfacial agents (i.e., clay intercalants) were additionally introduced during melt mixing on an open mill. Incorporation of hexadecyl trimethylammonium bromide (C16) resulted in three peaks in the XRD spectrum for both types of flocculants. The low basal spacing is due to Mt layers possessing inorganic cations. The medium basal spacing was assigned to Mt aggregates which were intercalated by C16 during melt blending via cation exchange. The highest basal spacing was attributed to Mt–C16 layers, which were further separated by intruded SBR chains. The interfacial agent 3-aminopropyltriethoxysilane (KH550), which is more polar than C16, yielded one peak in the XRD spectrum after melt blending. This cation exchange-induced peak was split in two peaks representing two clay populations after vulcanization of the acid-flocculated system. The Mt was always less intercalated by SBR molecules for the Ca2þ than for

PROPERTIES

387

the Hþ flocculated system [77]. Similar intercalated clay populations were produced when the intercalants were first incorporated in the aqueous suspension of Mt followed by the addition of the SBR latex [78]. In other cases, no peak could be resolved in the XRD spectrum for SBR/Mt vulcanizates produced via the latex method, at least up to 6 phr clay content. By further increase in the clay content (to 10 phr) a peak corresponding to a basal spacing of 1.37 nm appeared [79]. XRD spectra recorded at various treatment conditions for IIR/organoclay compounds delivered useful information on the role of heat and pressure during melt mixing. After melt blending of IIR with organically modified clay (OMC) on an open mill (without any curatives) the initial basal spacing of the OMC (2.24 nm) was shifted to higher values. By treatment at 80
C and atmospheric pressure, a reordering of the OMC platelets took place, which raised a peak corresponding to a basal spacing lower than that of the initial OMC [56]. This confinement may also be related with the alteration of the packing of the OMC intercalant. By increasing the treatment temperature the XRD reflections of the IIR/OMC changed in lesser extent. More specific, at 160
C, the shoulder at 5.88 nm in the XRD spectrum of the compound stored at 80
C disappeared. An intercalation peak around 2 nm was early raised, that is, after 10 min of treatment at 120
C and atmospheric pressure [56]. In case of solution mixed SBR with OMC, the increasing temperature of the heat treatment was accompanied with a lowered peak intensity in the XRD spectrum. This was attributed to a reduced degree of coherent layer silicate stacks and to the onset of further exfoliation [23]. On the other hand, when the pressure was increased to 15 MPa, heat treatment at higher temperatures resulted in a basal spacing of 3.15 nm instead of 2.08 nm measured after the heat treatment at 80
C (Figure 12.11). The gradual increase of the pressure at 160
C resulted in an increase of the peak intensity of the (001) reflection [23]. Peak absence in the XRD spectrum of IIR/organoclay vulcanizate, suggesting exfoliated morphologies, has been reported for IIR-g-MA [57]. Exfoliated organoclay structures, based on XRD characterization, were reported for NR/ENR systems [62].

12.5 PROPERTIES The tremendous interest and vivid research in diene rubber–clay nanocomposites is owing to their enhanced performance compared to conventional fillers, like carbon black and silica, at the same filler fraction. Due to the high aspect ratio of welldispersed clay platelets in the matrix the related composites exhibit very interesting properties, which are well manifested in their mechanical, wear, barrier, and fire resistance properties. 12.5.1 Mechanical (Dynamic–Mechanical) The peculiarities of clay, related to its structure, have been early recognized and exploited in rubber mixes. The alignment of the anisotropic clay particles in the direction of calendering resulted in a higher effective volume loading in that

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9 87 6 5

4

d (nm)

3

1

2

2400 * 3.15 nm

2100

Intensity (counts)

1800

15 MPa × 160ºC ×1 h

3.10 nm * * 2.11 nm

1500 2.75 nm *

1200

15 MPa × 140ºC ×1 h * 2.10 nm

900

15 MPa × 100ºC ×1 h 2.08 nm *

600 15 MPa × 80ºC ×1 h

3.50 nm *

300

Untreated

0 1

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3

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5

6

7

8

9

10

2θ (º)

FIGURE 12.11 WAXD patterns of the solution-intercalated IIR/OMC systems before (untreated) and after heat treatments. Notes: heat treatments done at different temperatures for 1 h at 15 MPa constant pressure. The asterisks indicate the (001) peak of the OMC dispersed in IIR matrix. The dotted line shows the location of the (001) reflection peak of the pure OMC. The curves were shifted vertically for clarity [23]. Copyright Wiley-VCH Verlag GmbH & Co. KGaA. Reproduced with permission.

direction [80]. Moreover, it has been well documented for conventional fillers, that with increasing surface area of the filler at constant loading the Payne effect becomes more pronounced. Additionally, the larger interfacial area in the unit volume compound at the same filler loading, results in more immobilized rubber shell. This increases the effective volume of the filler loading [81]. Therefore, the nanoscale dispersion of high aspect ratio platelets, suitably surface-modified or not, in a rubber matrix is expected to enhance both the filler’s surface area and the effective filler fraction. This should be accompanied with improved mechanical performance [82]. For example, NR vulcanizates containing 10 phr Mt-ODA showed similar mechanical properties as those reinforced with 40 phr carbon black (CB). Moreover, the NR/Mt-ODA exhibited almost two times higher tensile elongation at break value than the NR/CB [26]. The fine dispersion of organoclay platelets in the NR matrix is well reflected in the mechanical performance. NR reinforced with 5 phr Mt modified by aromatic intercalant presented tensile strength similar to the neat matrix. On the other hand, when NR was reinforced with 5 phr Mt modified by aliphatic intercalant, a threefold increase in tensile strength was measured [28]. In many trials the organoclay content of the related recipe was less than 10 phr. For NR mixed with Mt modified with octadecyl trimethylammonium bromide the maximum stress and strain at break values were recorded at a filler loading of 10 phr [83]. It has been shown that NR–organoclay nanocomposites, produced by solution mixing, yielded enhanced ultimate tensile properties and dynamic mechanical

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response compared to melt blended ones. This was attributed to higher bound rubber values in the compounds prepared by solution mixing [84]. The tensile properties are improved already at low clay contents [85,86]. However, the mechanical performance does not always deteriorate with higher clay loadings. For NR mixed with Mt via the latex route a significant enhancement of the modulus values was observed up to 30 phr Mt content. On the other hand, this was accompanied with a reduction of the elongation at break values [9]. Improved tensile properties due to clay dispersion are usually associated with enhanced tear strength, hardness, rebound resilience, as well as, with a decrease in the compression set, compared to conventional filled compounds for the same filler loading [25,66]. The NR compound with Mt-ODA exhibited additionally a remarkable lower heat buildup compared to NR with 50 phr CB [87]. Increasing further the clay loading (up to 70 phr) yielded pronounced reduction in some properties, such as rebound resilience, resistance to crack growth, and toughness [88]. The strain-induced crystallization of NR during a tensile test is well reflected in a Mooney–Rivlin plot (reduced stress against the reciprocal of the extension ratio). The abrupt upturn of the reduced stress at high extension in the related Mooney–Rivlin plot occurs at lower deformation for NR reinforced with organoclay compared to the neat NR vulcanizate [16,27]. This can be attributed to the limited extensibility of the NR chains, which bridge neighboring clay platelets and to enhanced rubber–filler interactions. Entanglements appear in the NR matrix due to the nanodispersion of clay platelets at low filler loadings. They may act as effective cross-links, which accelerate the crystallization process. The aspect ratio of the clay platelets has been found to play an important role on the onset of the first crystallites and the final crystallinity during the deformation of NR reinforced with organoclay. Organomodified Mt lent almost double crystallinity index at a specific extension in an NR matrix than organomodified laponite of lower aspect ratio. This is mainly related to the physical network, created in NR due to O-Mt filler, which also affected the morphology of the NR crystallites formed under elongation [75]. The filler-network created in the rubber matrix can be detected by the plot of dynamic shear modulus as a function of the shear strain amplitude. In this plot the region due to strong filler–filler interactions is well discernible [89,90]. For NR reinforced with 20 phr of organoclay an enhanced Payne effect is produced. As the strain increases the filler–filler “linkages” are broken [72]. In a double logarithmic plot of the excess modulus (measured at low strain in the linear stress–strain regime) against the filler fraction, the related curve scaled linearly with increasing the filler content. At higher filler loadings a steep slope increase was observed. This line intersection was traced to the reinforcing percolation threshold of the clay in the rubber matrix [91]. For NR reinforced with organoclay, the slope changed from 0.9 to 2.5 presenting a percolation threshold at 8 phr organoclay [92]. The percolation threshold can be assigned also to that clay loading below which the modulus of the vulcanizate is rather independent from the testing temperature (Figure 12.12). Above that point, the experiments carried out at lower temperatures deliver higher moduli than those at higher temperatures. Testing in the temperature range around 50
C is associated with a reduction of the efficiency of load transfer from NR to the clay. This

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

Initial modulus E (MPa)

20

Temperature (ºC) 25

16 12

Percolation threshold 70

8

100

4 0 0

0.01

0.02

0.03

0.04

0.05

0.06

True clay volume fraction

FIGURE 12.12 Initial modulus as a function of the true clay content for the sulfur-cured NR–organoclay samples (recipe without resorcinol), as obtained from tensile tests performed at cross-head rate of 100 mm/min and at temperatures of 25, 70, and 100
C [92]. Copyright 2009, with permission from Elsevier.

is due to a thermal transition of the intercalant of the organoclay. Thus, the immobilized NR chains in the vicinity of the clay platelets gain mobility as the alkyl tails of modifier (intercalant) transit to a disorder-like state [93]. In that temperature range the intercalated NR chains in the galleries of organo-Mt are surmised to trigger the onset of multiple endothermic peaks in DSC scans [75]. The rubber–filler interaction can be also detected through the plot of the mechanical damping against temperature [94]. The peak of the mechanical loss factor (tan d) was shifted to higher temperatures when organo-Mt served as filler in NR vulcanizates [75]. Moreover, the tan d versus temperature curve of NR/Mt-ODA presented the glass transition temperature (Tg) and also a second relaxation peak, the onset of which was assigned to intercalated/immobilized NR [66]. Reinforcing of BR with 3 phr organoclay resulted in a better mechanical performance compared to the same rubber when it was reinforced by 10 phr of CB [39]. Regarding the tensile and tear strength data as a function of the organoclay content, a property enhancement was detected up to 12 phr above which property deterioration started [17]. In other BR–organoclay nanocomposites, the property enhancement regarding the ultimate tensile properties was leveled off at 40 phr organoclay loading [37]. The role of the intercalant type was addressed for BR reinforced with 30 phr Mt modified by primary (P-OMt) or quaternary (Q-OMt) amine intercalants. P-OMt resulted in higher elongation at break values, whereas the Q-OMt yielded higher tensile and tear strength values, as well as, higher hardness and modulus. This was traced to a stronger stress-softening effect for BR/Q-OMt measured by cyclic tensile experiment, and a better interfacial adhesion was evaluated by a contact angle testing [38]. On the other hand, among several modified Mts, the one treated with dimethyl dihydrogenated tallow quaternary ammonium salt,

PROPERTIES

391

yielded the best mechanical performance in a BR matrix [40]. Nevertheless, additional incorporation of CTB compatibilizer (in 30 phr) in the related BR mixed with 10 phr Q-OMt enhanced the mechanical performance in both tear and tensile tests [41]. Adding organoclay in SBR improved the mechanical performance up to 10 phr organoclay loading [44]. This improvement diminished, however, with further increasing of the organoclay content [79]. Additional tensile property improvement was registered when the organoclay was pretreated with stearic acid in order to favor further clay intercalation [58]. At a filler fraction of 20 phr, a notable improvement in the mechanical behavior of SBR/Mt modified by KH550 was noticed when Si69 compatibilizer was also added in the compound. In that case, the Tg shifted by almost 15
C toward higher temperatures due to stronger rubber–filler interactions [51]. The mechanical performance of filled SBR vulcanizates was improved by the common use of organoclay and CB. The hybrid system containing 20 phr organoclay and 20 phr CB showed significantly increased tensile strength, storage modulus, and decreased damping performance due to the synergistic effect of the fillers [45]. It should be mentioned that the reduced intensity of the Tg peak suggests favorable interactions between the organoclay and the rubber matrix (in the cited case SBR) [48]. Property enhancement by hybrid fillers was reported also for an SBR solution blended with 5 phr sepiolite and 5 phr silica [95]. Furthermore, trials were made to use additional compatibilizers in rubbers with hybrid reinforcement (silica and organoclay). Simultaneous use of 10 phr of organomodified FHT, 20 phr silica and 3 phr bis(triethoxysilylpropyl)-tetrasulfan (TESPT or Si69) resulted in a significant increase in the tensile stress and strain at break values [73]. The related property improvement originates mainly from an extended clay platelet network that was generated in the SBR matrix. These platelets have been found to be fully oriented at about 300% strain during tensile test [53]. The percolation threshold in a double logarithmic plot of the excess modulus versus the filler fraction was recorded for SBR reinforced with organoclay. The threshold value was shifted toward lower filler fractions compared to conventional silica filled SBR vulcanizates [47]. Above the detected filler percolation threshold, the excess modulus, which is scaled according to a power law against the filler volume fraction, received an exponent of 3.9 for silica and 2.5 for the organoclay. This was related to the structural differences of the related filler networks [47]. It is noteworthy that by adding high amount of organoclay (up to 60 phr) a considerable amount of clay intercalant is introduced in the compound. This was well reflected in the dynamic mechanical analysis of the related SBR mixes [46]. Recall that in this example the intercalant content of organoclays lies at about 30 wt.%. On the other hand, when SBR was mixed with unmodified Mt via the latex route a significant improvement in both tensile and tear strengths was obtained even up to 20 phr clay loading [11]. Further clay loading did not deliver additional improvement in the mechanical properties [49]. For modeling the tensile mechanical performance of SBR–clay nanocomposites, the well-known Halpin–Tsai and Guth equations modified with a modulus reduction factor have been proposed [96]. The peculiar clay platelet network created by latex compounding in the SBR matrix was beneficial in

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

respect to other mechanical properties, which are usually less investigated. For example, a notable improvement was found for the flex-fatigue life of SBR vulcanizates reinforced with 5 phr clay and 25 phr CB [97]. The mechanical performance of IIR reinforced with a swollen O-Mt by butyl alcohol was enhanced in terms of tensile and tear strengths up to a filler content of 7 phr [98]. At higher O-Mt contents (up to 20 phr), the nanocomposites produced via solution compounding performed better than those prepared through melt blending. This was attributed to the higher aspect ratio of the organoclay in the vulcanizates delivered by the first method compared to the latter one [54]. A suitable way to improve further the mechanical performance of IIR–organoclay nanocomposites is to use IIR-g-MA instead of IIR [57]. Effects of the common incorporation of CB and organoclay in a chlorobutyl rubber (CIIR) matrix were investigated by response surface methodology. The statistical design of experiments covered the assessment of the mechanical properties and bound rubber content as a function of both filler loadings. The maximum tensile strength for the related vulcanizate was estimated for a recipe containing 27–32 phr CB along with 10–12 phr organoclay [99]. The superior action of clay to reinforce diene rubbers appears in their blends, too. It has been presented that incorporation of 10 phr ENR50 in NR followed by addition of 2 phr organoclay increases both strain and stress at break values compared to the nonreinforced blend [59]. Moreover, a second relaxation peak appeared in this compound at around 0
C that was assigned to intercalated and immobilized rubber fraction. The choice of 2 phr organoclay loading was based on the fact that at this clay content a maximum in the tensile and tear strength values was observed [61]. For NR mixed with SBR in 50/50 wt.% via the latex route, followed by addition of clay, an improvement in the ultimate tensile properties was detected up to about 5 phr clay content. Above this threshold the properties were leveled off [13]. 12.5.2 Friction/Wear/Abrasion The mechanisms involved in friction and wear of filled rubbers became subjects of extended research during the past decades. There are many different set-ups to investigate the friction and wear performances of rubbers. Moving a blunt needle on the rubber surface produces abrasive wear and the related surface track shows characteristic “waves of detachment.” In another case, as a hard sphere rolls over a rubber surface, the latter experiences along the rolling path a compression followed by relaxation. In this stress cycle, the energy is lost mostly due to hysteresis [100,101]. It is usually claimed that the abrasion wear of rubber correlates with the ultimate properties (i.e., tensile strength, elongation at break, tear resistance), whereas in respect to sliding wear, their fatigue and dynamic mechanical properties are of great importance [102]. It has to be born in mind that tribological properties are always system properties, and thus practically no correlations exist between friction/wear properties and structural/ mechanical properties. Nevertheless, some findings of the related research seem to be useful to develop organoclay-reinforced rubbers for tribological applications. It was proposed that the orientation of clay platelets in the vulcanizate may strongly affect the

PROPERTIES

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sliding wear behavior [103]. When fatigue crack growth is involved in failure under wear conditions, cavitations caused by decohesion between zinc oxide particles and rubber matrix may also affect the wear performance [104]. This phenomenon may be suppressed in rubber–clay nanocomposites [76]. Using a spherical steel intender of 3 mm in diameter that was reciprocating over a 5 mm long friction path under a load of 0.5 N, the friction properties of NR reinforced with organoclay were examined [87]. On an abrasion tester, NR compound containing 47 phr CB and 3 phr organoclay outperformed the reference compound, which contained 50 phr CB, in respect to the wear volume [105]. Lower abrasion loss of about 25% has been also shown for NR filled with 10 phr organoclay compared to the neat vulcanizate. This nanocomposite compound delivered the highest hardness and tensile strength [26]. On the contrary, an increase in wear has been found for NR–clay nanocomposites prepared via latex compounding [106]. The damping characteristics of NR reinforced with 10 phr organoclay at 0 and 60
C, respectively, have been used to speculate about the tribological performance of the related vulcanizate. An almost 5% higher tan d at 0
C along with 5% lower tan d at 60
C were measured for the nanocomposite compared to the plain NR vulcanizate. The first favors the wet grip of tires while the latter relieves the rolling resistance. The related compound exhibited increased tear and tensile strengths, as well as, hardness compared to the neat matrix [107]. The volumetric loss of cylindrical BR vulcanizates after their contact with the abrasive surface of a rotating drum under the load of 1 N has been evaluated. The BR–organoclay nanocomposite presented 60% higher abrasion resistance than the BR/Mt compound [40]. Attempts were also made to exploit the masterbatch technique to produce abrasion resistance rubber nanocomposites. ENR mixed with organoclay in 1 : 1 ratio, has been compounded in 10 phr in a blend of NR with SBR of high styrene content in a ratio of 80/20. When the reinforcement involved 40 phr CB combined with 5 phr organoclay, the abrasion resistance enhanced almost by 30% compared to the vulcanizate reinforced by 40 phr CB. The improvement was a function of the CB type used [21]. The wear performance of these vulcanizates has been measured on a circular rubber disc using different types of rocks as abraders. The rubber disc was turning at 246 rpm (i.e., 0.8 m/s linear velocity for the specific set-up). Applying a force of 4.4 N the wear loss of the hybrid-reinforced systems was markedly reduced. The compound that delivered the best tribological performance also showed the lowest compression set, highest tensile strength, and hardness among the specimens examined [21]. In another example, NR/BR compounds in 60/40 weight ratio, containing various amount of organoclay have been prepared. The filler fraction increase up to 15 phr yielded improved abrasion resistance. This was accompanied by a parallel enhancement in the tensile, tear strength, and hardness values [108]. 12.5.3 Barrier The property, which has been presented in the open literature to be strongly improved due to the platy structure of clay, is by far the barrier. The permeability of rubber–clay

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

nanocomposites is decreased by increasing clay content [109]. Furthermore, the platelet aspect ratio, the orientation and the degree of layered silicate exfoliation all affect the barrier performance of the related nanocomposites [110]. NR reinforced with 3 phr organoclay resulted in a 50% reduction of the oxygen permeability compared to the neat matrix. The same nanocomposite presented toluene absorption at 20
C almost by 40% lower than the pure NR vulcanizate. It has to be mentioned that the resistance to toluene swelling appeared already by adding 1 phr organoclay (30% reduction compared to the plain NR). Up to 5 phr organoclay still some reduction in the toluene absorption was noticed [86]. In another series of NR–organoclay nanocomposites, addition of 5 and 10 phr organoclay resulted in about 10 and 15%, respectively, reduction of the oxygen permeability compared to the plain NR vulcanizate. NR/10 phr organoclay mix showed a 30% reduction for toluene absorption at 30
C compared to the pure NR vulcanizate [111]. It seems that the gas permeability of organoclay reinforced NR is in accordance with the solvent swelling behavior. For NR–clay nanocomposites, prepared via the latex route, the barrier properties were enhanced even at lower clay content. Addition of 1 phr clay in NR delivered more than 35% decrease in the oxygen permeability. A further clay increase for 2 and 3 phr reduced the oxygen permeability by about 45 and 50%, respectively, compared to the unreinforced matrix. The same series of specimens presented a lower calculated diffusion coefficient for toluene at 30
C than the pure NR (15% less for 1 phr clay and 40% less for both 2 and 3 phr clay contents). On the other hand, the absorption coefficient in toluene did not follow the same trend as the diffusion coefficient [112]. It is most likely that the labyrinth structure created by the clay platelets has a greater effect on the solvent diffusion than on the solvent uptake at equilibrium. The labyrinth effect (tortuous path) has been investigated as a function of the clay aspect ratio for NR/clay vulcanizates. The clay was introduced in NR by latex compounding. It has been shown for NR with 10 phr clay content that FHT yielded almost 40% lower, whereas Mt about 30% lower toluene uptake at 25
C compared to the neat matrix [8]. Furthermore, the initial circular specimens, cut for the swelling measurements, changed their shape differently due to toluene immersion. NR gave the highest diameter and the lowest thickness increase, whereas NR/FHT vulcanizate registered the highest thickness along with the lowest diameter increase. The orthotropic dispersion of clay platelets triggered specific dimensional changes, which were amplified by the clay orientation and clay aspect ratio. For highly filled unvulcanized NR–Mt nanocomposites prepared by latex compounding, the xylene uptake, which was already reduced at 5 phr Mt, was further strongly reduced when the Mt content reached 30 phr. The latter compounds showed the highest dimensional anisotropy upon swelling toluene uptake at 25
C [9]. The oxygen permeation behavior of hybrid reinforced (calcinated clay, silica, and CB), melt-blended NR composites has also been studied. Incorporation of CB in the NR/clay/silica vulcanizate resulted in an increase of the oxygen permeability as the primary size of CB particle increased [113]. It was speculated that the rubber–filler interactions play an important role in the gas permeation of rubber nanocomposites [114].

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Addition of organoclay to BR resulted in about 20% lower toluene uptake at equilibrium, compared to BR matrix, as the amount of clay increased up to 10 phr [17]. On the other hand, the water vapor permeability of plain BR vulcanizate was reduced by almost 80% when 5 phr organoclay was introduced. Incorporation of compatibilizer did not improve this excellent barrier performance further [41]. SBR/organoclay masterbatch has been produced by modifying clay with ODA during the latex mixing procedure in a sodium chloride rich suspension. The SBR/ organoclay compound, produced by incorporation of the above masterbatch, yielded an almost 15% reduction in toluene swelling compared to SBR melt blended with clay at the same content, that is, 5 phr [115]. In another case, the gas permeation of SBR, latex blended with 20 phr Mt, has been measured. It was found that the related nanocomposites exhibited a 55% reduced nitrogen permeation compared to the neat vulcanizate [11]. With increasing amount of the clay in SBR, the nitrogen permeability was strongly reduced. For SBR, latex blended with 40 phr rectorite, the permeation of nitrogen was more than 70% lower than that of the pure matrix [116]. The nitrogen permeability of highly filled SBR/organoclay has also been investigated (Figure 12.13). When SBR was melt blended with 150 phr Mt-ODA the nitrogen permeability was reduced by about 85% compared to the plain matrix [46]. The barrier behavior of SBR–clay nanocomposites has been investigated for hybrid filler-reinforced systems, too. By solution blending, SBR was mixed with 5 phr sepiolite and 5 phr silica in presence of Si69 compatibilizer. The hybrid fillerreinforced mix exhibited about 55 and 30% lower swelling in toluene compared to the neat matrix and to the SBR filled with 5 phr sepiolite, respectively [95]. On the contrary, incorporation of 5 phr Mt together with 25 phr CB in SBR matrix via latex blending did not register any improvement in toluene sorption behavior [97]. 1,0

0,8

Pc /Pp

0,6

0,4

0,2

0,0 0

20

40

60 80 100 120 Clay loading (phr)

140

160

FIGURE 12.13 The influence of organoclay loading on N2 gas permeability in an SBR stock. Note that Pc and Pp represent the gas permeability of the composite and the pure matrix, respectively.

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

The good barrier property of IIR has been further improved by addition of clay. In an example, IIR-g-MA was melt blended with Mt-ODA. Incorporation of 15 phr organoclay resulted in about 60% reduction of the nitrogen permeability [57]. The tortuous path, created in the IIR matrix due to the organoclay dispersion, was found to be affected by the curing temperature. For a clay intercalant, which was able to react with the sulfur curatives at high temperatures, vulcanization above 150
C caused a reaggregation of clay platelets. This process lowered the aspect ratio of the clay platelet stacks at the same filler fraction that supported the nitrogen molecules to pass through the rubber matrix. Accordingly, curing at 180
C yielded about 30% gas permeation increase than curing of the same compound at 150
C [55]. To maintain the high aspect ratio of the clay platelets, IIR was mixed with organoclay through a solution technique. Incorporation of 5 phr organoclay in IIR via solution blending resulted in an almost 10% decrease in nitrogen permeability compared to IIR/ organoclay of the same filler content produced by melt blending [54]. In order to enhance further the barrier property of IIR–clay nanocomposites, the organoclay was first swollen in butyl alcohol, followed by melt blending with IIR. The resulted nanocomposite with 5 phr filler presented about 5 and 20% lower nitrogen permeability compared to a solution mixed IIR–organoclay nanocomposite and the IIR vulcanizate, respectively [98]. The barrier performance has been investigated also for various diene rubber blends. The swelling index in toluene at room temperature of NR blended with 10 phr ENR decreased by almost 10% when the related stock contained 2 phr organoclay [60]. The combined action of ENR and organoclay to improve the barrier properties has been verified using a blend of NR/SBR (80/20) reinforced with 40 phr CB. The compatibilization of the related blend by an ENR masterbatch bearing 10 phr organoclay resulted in enhanced solvent and oil resistance at 25 or 100
C [21]. Likewise, ENR, solution-mixed with 50 phr clay, was exploited as compatibilizer in NBR. By adding 10 phr ENR masterbatch, the oil and solvent uptakes of the corresponding NBR formulations were reduced [117]. In another case, Mt itself acted as compatibilizer for SBR/NBR compounds. Incorporation of Mt in the related blends was accompanied with markedly reduced swelling in toluene. More specific, the addition of 20 phr Mt in SBR/NBR (50/50) resulted in 36% less swelling compared to the neat vulcanizate [118]. 12.5.4 Fire Resistance In demanding applications, diene rubber articles are subjected to severe conditions also including contacts to flame and fire. The ongoing development with fireresistance diene rubber compounds, especially requested by the aviation industry, was extended also for the incorporation of organoclays in rubbers. This research direction was reasoned by the success achieved with thermoplastic/organoclay systems in this field. NR reinforced with Mt modified by tributyl phosphate (TMt) was tested by cone calorimetry at an incident heat of 35 kW/m2 [119]. Addition of unmodified Mt in NR yielded a reduction of the heat release rate (HRR), especially when the Mt content

PROPERTIES

397

Heat release rate (a.u.)

Rubber/microfiller Pure rubber

Rubber/nanofiller

Time (a.u.)

FIGURE 12.14 Comparison of the heat release rate (HRR) versus time plots for pure rubber and its micro- and nanocomposites schematically.

increased up to 20 phr. By incorporating TMt in the mix the HRR was further reduced. Increasing amount of TMt (up to 30 phr) delivered an almost 40% lower HRR than the neat NR matrix. The mass loss rate also decreased for the related organoclay reinforced compounds. On the other hand, the value of smoke produce rate presented a similar decrease for all NR–clay micro- and nanocomposites. This performance may be linked with the char creation on the surface of the vulcanizates. Flammability tests were performed on SBR/Mt compounds produced via the latex route. The curing agents were incorporated in the related mixes on an open mill. Using a cone calorimetry under a heat flux of 50 kW/m2 the SBR filled with 20 phr Mt produced 27% lower HRR peak than the pure matrix. Additionally, the nanocomposite delivered the longest time to ignite, the lowest mass loss rate, as well as, the largest amount of char upon combustion compared with the microcomposite and the neat matrix [120]. The effect of the filler type on the heat release rate versus time traces is summarized schematically in Figure 12.14. 12.5.5 Others Apart from the widely known and popular properties of diene rubber–clay composites, the platy structure of nanodispersed clay in the rubber matrix may affect other less cited properties. NR latex mixed with Mt along with various coupling agents was found to alter the ozone aging resistance. Aging experiments of the specimens, which were put additionally under 20% strain, were carried out in an ozone oven. The measured ozone cracking time of NR reinforced by 20 phr clay generally decreased by incorporation of coupling agents or antioxidants. On the other hand, the resistance to ozone attack of NR reinforced with 20 phr clay could be increased by more than

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RUBBER–CLAY NANOCOMPOSITES BASED ON APOLAR DIENE RUBBER

10% by the combined action of an amine antioxidant and a titanate coupling agent [121]. Due to clay exfoliation, an initially transparent rubber matrix may maintain this property also in the corresponding rubber–clay nanocomposites. For BR reinforced with an organoclay fraction lower than 5 phr, the haze of the peroxide cured compounds was not affected. By contrast, with further increasing filler loading the haze was grown gradually, mainly due to clay aggregation phenomena [37].

12.6 APPLICATIONS AND FUTURE TRENDS Considering clay as a substitute of conventional fillers, like CB and silica, a significant weight reduction comes directly into forth. Recall that mechanical property improvement of diene rubber/clay vulcanizates can be achieved just at low filler contents. At the same time, the nanoscale-dispersed clay does not influence the ultimate elongation values. Therefore, the related nanocomposites may find applications in all sectors where rubber products are traditionally used. The low heat build-up combined with enhanced wear resistance is promising features for tires and technical rubber goods subjected to wear (sealing, bushing, etc.). Taking additionally into account the notable barrier improvement due to clay platelets, nanocomposites may find applications in barrier sheets, inflatable tubes, membranes, gloves, tread and innerliner of tires, and so on. The enhanced nitrogen barrier performance delivered by diene rubber–clay nanocomposites may be exploited in demanding articles such as tires for racing cars and aircrafts, where nitrogen mixtures are used for inflation and the related parts exposed to high pressure and a broad temperature range. The enhanced abrasion resistance of diene rubber–clay nanocomposites may be exploited in coating applications, even when fire resistance is a requirement. The coloring ability of diene rubber–clay nanocomposites along with their enhanced mechanical performance may assist their penetration in articles of everyday life (household, garments, footwear sector, etc.). Since diene rubber–clay nanocomposites meet the requirements of suitable rebound resilience, durability and traction, along with nonmarking performance, shoe sole applications are also of relevance [122]. Dispersion of clay in diene rubber vulcanizates may tune the acoustic impedance of the related articles. The dynamic behavior of diene rubber–clay nanocomposites along with their moderate density may be suitable for navy applications, like acoustic tiles in submarines, sonar domes, and the like [123]. However, it has to born in mind that the properties of diene rubber–clay nanocomposite highly depend on their manufacturing process. For latex blending, issues related to coagulation should always be taken into account since it is the source of clay aggregation [124]. On the other hand, the choice of a suitable solvent (by considering the solubility and the interaction parameters) for the solution mixing technique may be beneficial to achieve the optimum performance of the related rubber nanocomposites [125]. In melt compounding procedures, it is always desirable to monitor the progress of mixing. For diene rubber/organoclay mixtures, the

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appropriate organoclay dispersion in the matrix generally determines the final properties of the vulcanizate. Online measurements of electrical conductance may be used for detecting the degree of clay platelet dispersion in the rubber matrix during melt blending [126]. Additionally, optimization of the curing recipes may further improve the properties of the vulcanizates. For example, homogeneous dispersion of ZnO particles in a diene rubber matrix will minimize the cavities created by rubber/ ZnO decohesion, and thus enhancing the fatigue resistance [104]. Recipes should be examined on how to reduce the presence of Zn þ2 ions providing more ecological friendly compounds [127]. Moreover, suitable coupling agents and ternary hybrid filler systems, involving CB and silica, may also enhance the performance of diene rubber–clay nanocomposites [128]. For example, additional generation of starch nanocrystals in rubber–clay nanocomposites may improve their property profile and render them more eco-friendly [129]. A method that has not been extensively followed in the open literature for diene rubber–clay nanocomposites, although it is able to produce exfoliated clay structures, is the in situ polymerization technique [130].

ACKNOWLEDGMENT This work is connected to the scientific program of the “Development of qualityoriented and harmonized Rþ D þI strategy and functional model at BME” project. This project is supported by the New Hungary Development Plan (Project ID: ´ MOP-4.2.1/B-09/1/KMR-2010-0002). TA

REFERENCES 1. Sommer, F. Aufbau, Herstellung und Eigenschaften von Kautschuk und Elastomeren. In Kautschuktechnologie, F. R€othemeyer, F. Sommer (eds.), Hanser, Munich, 2001, pp. 41–149. 2. Hofmann, W. Rubber Technology Handbook, Hanser, Munich, 1994, pp. 230–231. 3. Theng, B. K. G. Formation and Properties of Clay–Polymer Complexes, Elsevier, Amsterdam, 1979, pp. 3–36. 4. Van Olphen, H. An Introduction to Clay Colloid Chemistry, Wiley, New York, 1977. 5. Vaia, R. A. Structural characterization of polymer-layered silicate nanocomposites. In Polymer–Clay Nanocomposites, T. J. Pinnavaia, G. W., Beall (eds.), Wiley, Chichester, 2000, pp. 229–266. 6. Shen, S.; Yang, M.; Ran, S.; Xu, F.; Wang, Z. Preparation and properties of natural rubber/palygorskite composites by co-coagulating rubber latex and clay aqueous suspension. J. Polym. Res., 13, 469–473 (2006). 7. Karger-Kocsis, J.; Wu, C. M. Thermoset rubber/layered silicate nanocomposites. Status and future trends. Polym. Eng. Sci., 44, 1083–1093 (2004). 8. Varghese, S.; Karger-Kocsis, J. Natural rubber-based nanocomposites by latex compounding with layered silicates. Polymer, 44, 4921–4927 (2003).

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CHAPTER 13

RUBBER–CLAY NANOCOMPOSITES BASED ON NITRILE RUBBER KONSTANTINOS G. GATOS JÓZSEF KARGER-KOCSIS

13.1 INTRODUCTION The driving force to copolymerize 1,3-butadiene and acrylonitrile (AN) at the beginning of 1930 was the quest for solvent resistance elastomers. The related rubber, widely known as nitrile rubber (NBR derived from nitrile–butadiene rubber) was initially commercialized in Europe and United States by I.G. Farbenindustrie and B.F. Goodrich, respectively [1]. Nowadays, several manufacturers produce NBR versions worldwide including Nippon Zeon, Uniroyal Chemical, and Lanxess [2]. NBR is an amorphous elastomer, its structure may vary from branched to linear one depending on the conditions of polymerization. In terms of polymerization temperature, NBR produced at lower than 25
C is classified as “cold NBR” compared to its “hot” version polymerized at higher temperatures. The AN content of NBR types is usually between 15 and 53 wt.% [3]. High AN content ensures excellent oil resistance, however, the cold flexibility of the corresponding NBR is sacrificed. NBR applications involve seals, O-rings, timing belts, hose tubes, oil-drilling parts, gaskets, gloves, and footwear. In order to improve the heat, weather, and oxidation resistance of NBR, its unsaturation (i.e., double bond content) should be reduced. This resulted in the development of hydrogenated nitrile rubber (HNBR). For increased abrasion resistance of nitrile rubber, copolymerization of methacrylic or acrylic acid with acrylonitrile and butadiene yields NBR grades containing pendant carboxylic groups (XNBR) [3]. The latter rubber might be considered as the most widely used NBR terpolymer. The polarity and properties of NBRs differ from those of other butadiene-like rubbers. This is the major reason why (H)NBR is separately treated within the family of the diene rubbers.

Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti. Ó 2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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Traditionally, nitrile rubber has been reinforced with carbon black or silica in order to enhance its mechanical performance [1]. Nevertheless, several types of fillers have been presented in the literature to enhance specific properties of this elastomer [4–6]. On the other hand, layered clay has received less attention for compounding with NBR. Its action was expected to follow the reinforcement ability of micron-scaled platy fillers [7,8]. In the beginning of 1990, it was shown that the clay can be efficiently and nanoscale dispersed in polymer matrices [9,10]. This triggered some early works on NBR–clay nanocomposites [11,12]. Most of the clay minerals belong to the category of layered silicates or phyllosilicates [13]. Under certain conditions macromolecular chains might be present in confined form within the silicate galleries. This expands the basal (interlayer) spacing of the layered silicate sheets (having an individual sheet thickness of about 1 nm). Depending on the fact whether the clay layers are still regularly stacked or dispersed in a random manner in the matrix the related systems are termed intercalated and exfoliated nanocomposites, respectively [14]. Apart from the widely used smectitetype silicates, which belong to the family of cationic clays, layered double hydroxides, representing anionic-type clays, are also available and used for nanocomposite preparation [15]. In order to expand the interlayer distance of the layered clays the initial cations or anions are replaced by low molecular weight surfactants of cationic or anionic nature. This renders the clay organophilic and thus more compatible with the polymer than the untreated (pristine) one. The outcome of this organophilic treatment (i.e., clay modification via ion exchange reaction) is usually quoted as “organoclay.” NBR has been successfully reinforced with layered clay during the last decade and several factors have been identified to affect the final properties of the related vulcanizates. The research conducted on this field revealed the great effect of the compounding methods on the NBR–clay nanocomposite formation.

13.2 PREPARATION METHODS AND CLAY DISPERSION 13.2.1 Solution By dissolving a rubber in a suitable solvent and introducing a dispersion of organoclay in an appropriate solvent in it, the clay platelets may be nanoscale dispersed. When the solvent is left to evaporate, the clay layers might remain evenly exfoliated in the elastomer matrix. It should be mentioned that the selected curatives are introduced either directly in the solution or after the solvent evaporation step through conventional compounding techniques [16]. Hwang et al. [17] dissolved NBR with 32 wt.% AN content in methyl ethyl ketone (MEK). The organoclay dispersed in the same solvent was montmorillonite (Mt) modified with dimethyl dehydrogenated tallow quaternary ammonium salt. Both solutions were mixed and stirred for 12 h. After evaporating the MEK for 2 days under vacuum, sulfur-type curatives were added on an open mill. For organoclay loadings up to 10 parts per hundred parts rubber (phr) the related vulcanizates presented intercalated clay structures. This was confirmed by means of X-ray diffraction (XRD) scattering.

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In the respective XRD spectra, the peak corresponding to the basal spacing of the organoclay, was shifted to lower scattering angles (2u) notifying that rubber chains were intercalated within the interlayer of the organoclay. The largest interlayer spacing was found at 3 phr organoclay content. Nevertheless, transmission electron microscopy (TEM) signified also exfoliated organoclay platelets in the NBR matrix. Chloroform has also been used as solvent for NBR [18]. Sadhu and Bhowmick [19] dissolved NBR in chloroform wherein Mt modified with octadecyl ammonium chloride, dispersed in ethyl alcohol, was added. The required amount of organoclay dispersion was stirred with the rubber solution followed by incorporation of peroxide and the mixture was dried in an oven for 2 days at 50
C. At an organoclay loading of 4 phr, the NBR with 19 wt.% AN content produced apart from intercalated structures also exfoliated ones. With increasing AN content of the NBR (34 and 50 wt.%), no exfoliated structure was detected in the respective TEM images. Accordingly, the increased polarity of the matrix was accompanied with clay aggregation. 13.2.2 Latex Clay is hydrophilic and thus water/clay slurries can be easily prepared. In case of swellable layered clays (e.g., Mt, hectorite) the water molecules penetrate within the galleries due to hydration of the interlayer cations [13]. Thus, through stirring of the slurry the clay layers can be peeled off and individually dispersed in the water. As NBR exists also in latex form (being an aqueous dispersion of rubber nanoparticles), its mixing with clay in water medium is a straightforward technique to achieve nanocomposites. Wu et al. [20] stirred NBR latex of 26 wt.% AN content with the aqueous suspension of Mt followed by co-coagulation with dilute sulfuric acid. The suspension washed with water (till pH to become 7) and dried in an oven for 24 h at 70
C. Sulfuric curatives were added on a two-roll mill. The vulcanizates with 20 phr filler loading presented stacks of Mt platelets dispersed in the NBR matrix, as this was revealed through TEM images. Following a similar latex mixing route NBR/clay vulcanizates with silicate loadings of 10 and 20 phr, respectively, were produced. No difference was found in the 2u peak positions of the XRD spectra in function of the silicate content [21]. In order to improve the clay dispersion in the NBR matrix, Hwang et al. [22] adapted the ball milling technique. A mixture of Na-Mt along with dispersing agent (aromatic polyglycol ether), nonionic emulsifier, electrolyte (potassium hydroxide), and deionized water was prepared in a ball-milling equipment. This mixture was stirred for 48 h with NBR latex (AN content of 31–33 wt.%). Afterwards sulfuric curatives were incorporated, also via ball milling, and the mixture was stirred for additional 48 h. The slurry was left to dry for 24 h at 50
C prior to curing. As detected by XRD, the vulcanizates presented intercalated nanocomposite structures at all clay loadings (i.e., 3, 5, and 7.5 phr). More specifically, the basal spacing of Mt was expanded from 1.29 nm up to about 3.30 nm for the NBR/Mt vulcanizates. Moreover, TEM images displayed a number of exfoliated clay platelets especially for low filler content. The latex route has been followed by several research groups. Kader et al. [23] stirred NBR latex (AN content: 34 wt.%) with clay dispersion extensively for 3 h at

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RUBBER–CLAY NANOCOMPOSITES BASED ON NITRILE RUBBER

(a)

(b) Na-MMT

Intensity (a.u.)

7.37º (1.20 nm)

Na-MMT+DCA 6.27º (1.41 nm) NBR-10M 6.35º (1.39 nm) NBR-5M NBR-2M NBR

2

4

6 2θ (º)

8

10

FIGURE 13.1 (a) Wide-angle XRD for Na-Mt (pristine clay), NBR, and NBR–organo-Mt nanocomposites with varying proportions of organoclay. Designations: 2, 5, and 10M designate 2, 5, and 10 phr organoclay contents, respectively. Notes: MMT ¼ Mt, both Na-Mt and Na-Mt/DCA are also included for comparison. (b) High magnification TEM image of NBR–organo-Mt (5 phr) intercalated–exfoliated nanocomposite [23]. Copyright 2006, with permission from Springer Science þ Business Media.

60
C whereas the electrolyte, used for coagulation, was a dilute solution of dichloroacetic acid (DCA). The pH of the mixture was adjusted to 5–6 (the initial pH was 4) and the blend dried for 18 h at 80
C. Finally, sulfur-type curatives were added on an open mill. As shown in Figure 13.1, the XRD spectra of the respective vulcanizates presented apart from the broad scattering at low 2u angle (suggesting the presence of exfoliated clay platelets) also scattering around 6.35
(indicating intercalated clay platelets). However, such a small interlayer expansion of the silicate (from 1.20 to 1.40 nm) was ascribed to the effect of DCA. On the other hand, TEM investigation detected nicely exfoliated clay populations, as well. The special role of the additives incorporated during latex compounding was presented by Wu et al. [24]. NBR latex with 24–26 wt.% AN content was co-coagulated with clay using 1% calcium chloride aqueous solution followed by drying in an oven. It was found that during coagulation the initially exfoliated clay platelets in the solution are forced to reaggregate by the flocculant. The latter might be present within the clay galleries. Similar results were also obtained in case of XNBR bearing 31–35 wt.% AN content [24]. However, in presence of flocculant, the smaller the rubber latex particles, the higher the latex content and the faster the co-coagulation are the higher the chance is to receive nanocomposites with exfoliated clay. 13.2.3 Melt Blending Because melt compounding is the usual rubber mix formulation, this preparation method received most of researchers’ attention also for the production of NBR–clay nanocomposites. In an example, NBR bearing 34 wt.% AN content was blended with

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organoclay on an open two-roll mill followed by the incorporation of the sulfuric curatives [25]. Pretreatment of NBR with phenolic resin followed by the addition of 5 phr organoclay was found to favor intercalated nanocomposite structures of high interlayer spacing in the respective sulfur-cured vulcanizates [26]. In order to investigate the polarity effect, Liang et al. [27] mixed NBR of various AN contents (i.e., 26, 35, or 42 wt.%) with 10 phr organoclay for 15 min on an open mill. The increased polarity of NBR assisted marginally the dispersion during melt mixing, although the interlayer spacing of the organoclay rose with the NBR polarity in the related vulcanizates. Although open mill process is convenient for rubber compounding, it has been presented that nanocomposite formation might be favored by intensive mixing and selecting high compounding temperatures [28]. Such conditions support an efficient dispersion of the silicate layers and stacks in the rubber matrix. Accordingly, compounding NBR in an inner mixer with organoclay has delivered nanocomposite structures [29]. However, for that an adequate length of the alkyl tail of the clay intercalant was required. Kim et al. [30] showed for NBR having 29 wt.% AN content that the organoclay was exfoliated in the matrix when octadecylamine (ODA) or dodecylamine instead of octylamine served as the initial intercalant (surfactant) of the clay. All components were mixed for 30 min at 50 revolutions per minute (rpm) in an internal mixer at room temperature. Likewise, blending at higher temperatures has delivered exfoliated organoclay structure. When NBR (34 wt. % AN content) was mixed with organoclay in an inner mixer for 10 min at 90
C TEM inspection revealed well dispersed organoclay stacks [31]. Usually, a two-step melt-blending approach is adapted to prepare rubber–clay nanocomposites. This includes the mixing of rubber with organoclay in an inner mixer at high temperature followed by addition of curatives on an open mill at low temperature. Thus, desirable clay dispersion without activating the curatives can be reached. Das et al. [32] compounded NBR of 44 wt.% AN content with 5 phr Mt modified with a quaternary ammonium salt in an internal mixer for 10 min at 160
C and 50 rpm rotor speed. Incorporation of the curatives took place on an open mill for 10 min at 40
C. This mixing protocol yielded intercalated and exfoliated structures as evidenced by means of XRD and TEM, respectively. Conventional melt mixing techniques have been also exploited in case of HNBR–organoclay nanocomposites. HNBR has been compounded with organoclay on an open mill [33] or in an inner mixer [34]. In the latter case, the curatives were incorporated on a two-roll mill. For a HNBR of high AN content (i.e., 43 wt.%) when mixed with Mt that was modified with methyl-tallow-bis(2-hydroxyethyl) quaternary ammonium salt, the clay became efficiently intercalated. This was traced to the fact that this intercalant is able to create hydrogen bonds with the AN groups of HNBR [35,36]. Ali et al. [37] based on online electrical conductance measurements characterized the organoclay dispersion during melt blending in an inner mixer. Initially, HNBR of 33 wt.% AN content was compounded with peroxide at 50
C and 70 rpm rotor speed. After 7 min, clay modified with stearyl benzyl dimethyl ammonium chloride was incorporated in 5 phr loading. The online measured electrical conductance was found to increase after the addition of the organoclay and to reach a plateau after a certain time. This behavior was correlated with the

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RUBBER–CLAY NANOCOMPOSITES BASED ON NITRILE RUBBER

progress of the exfoliation of organoclay stacks within the rubber matrix, whereby observing a plateau region (passing the so-called “characteristic point”) an adequate clay dispersion degree was reached. Similar analysis was also performed for XNBR, which delayed the onset of the “characteristic point” compared to HNBR [37]. Melt compounding has been followed in case of XNBR mixed either with cationic [38] or anionic clays [39]. In both stocks blending of the matrix with the clays was performed in an internal mixer at 100
C followed by the addition of the curatives. Melt blending of an XNBR (with 7 wt.% carboxylic acid monomer content) with 5 phr layered double hydroxide (LDH) modified with dodecylbenzene sulfonate resulted in a torque increase during kneading. That was attributed to the interactions between the CN and COOH groups of XNBR with the OH functionalities of the modified LDH. The torque increase was further amplified when 5 phr of ZnO was added in the formulation [39]. It has been reported that incorporation of organoclay in NBR using epoxidized natural rubber (ENR)/organoclay masterbatch is beneficial. ENR was initially solution mixed with 50 phr organoclay followed by drying. The dry masterbatch was blended with NBR on an open mill along with sulfur-type curatives. This resulted in mainly exfoliated and intercalated dispersions at 5 and 10 phr organoclay contents, respectively [40].

13.3 CURE CHARACTERISTICS Vulcanization is considered as an important step in obtaining useful rubber articles. Taking into consideration the type of rubber and the intended application, the related vulcanization recipe and conditions are carefully selected. In common practice, NBR is cured either by sulfur or peroxide. In the latter case, vulcanizates with increased thermal stability are obtained [1]. The chemistry involved during curing of nitrile rubber stocks reinforced with modified layered silicates has raised vivid interest in academia. In presence of Mt modified with dimethyl dehydrogenated tallow quaternary ammonium salt, the scorch time of NBR, cured by sulfur, was reduced even at 1 phr organoclay loading. This was attributed to the acidic nature of the organosilicate that activated the formation of soluble zinc ions inducing the decomposition of the accelerator in an earlier stage [17]. Scorch time can be further depressed by increasing amount of the organoclay [41]. As detected by differential scanning calorimetry (DSC), the organoclay reduces the activation energy of vulcanization of NBR during sulfur curing compared to the nonmodified (pristine) clay reinforced mix [25]. Comparing the sulfur and peroxide curing for NBR having 44 wt.% AN content, Das et al. [42] investigated the intercalation degree of clay modified with distearyl dimethyl quaternary ammonium chloride. As shown in Figure 13.2, the organoclay was intercalated by NBR during melt compounding. The initial basal spacing of the organoclay at 2.98 nm (not shown here for brevity reasons) was shifted to 3.57 nm during mixing in a kneader for 10 min at 160
C and 50 rpm rotor speed. Incorporation of the curatives took place on an open mill for 10 min at 40
C, followed by

CURE CHARACTERISTICS

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Sulfur cured + 5 phr organoclay

(a) 40,000

d=0.72 nm

20,000 d=4.35 nm d=1.97 nm (001) (002) d=1.34 nm (003)

80,000 (b)

Peroxide cured + 5 phr organoclay

d=3.69 nm (001)

Intensity

60,000

40,000 d=1.88 nm (002)

20,000

d=1.28 nm (003)

d=1.00 nm (004)

50,000 (c)

Masterbatch + 5 phr organoclay

40,000

30,000

20,000

d=3.57 nm (001)

10,000

d=1.84 nm (002) d=1.26 nm (003)

0 0

2

4

8 6 2θ angle

10

12

14

FIGURE 13.2 WAXD patterns of NBR reinforced with 5 phr organoclay and cured by (a) sulfur, (b) peroxide, and (c) NBR masterbatch (without curatives and thus uncrosslinked) [42]. Copyright 2008, with permission from Wiley-VCH Verlag GmbH & Co. KGaA.

vulcanization at 160
C. It was found that both vulcanizates yielded intercalated nanocomposite structures whereas sulfur curing raised higher interlayer spacing (i.e., 4.35 nm) compared to the peroxide cured version (i.e., 3.69 nm). It seems that sulfuric ingredients favor the organoclay intercalation during melt blending [43].

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In order to increase further the degree of clay intercalation in sulfur-cured NBR, an excess amount of stearic acid (SA) was used. Adding 4 phr SA instead of 1 or 2 phr (whereby the other sulfur curatives remained the same) expanded the interlayer spacing of the organoclay to 4.70 nm compared to 4.12 (1 phr SA) and 4.35 nm (2 phr SA), respectively [42]. On the other hand in their sulfur-type recipe, Ma et al. [44] first pretreated the organoclay with SA (in a ratio of 10:2) in a mixer for 1 min, followed by thermal treatment in an oven at 100
C. This low molecular weight SA was found to intercalate within the clay galleries, assisting further the organoclay dispersion in the vulcanizate. Due to the fact that curatives are able to penetrate the clay interlayers, to consider their interaction with the organic surfactants of the clay is quite reasonable. Effects of various clay intercalants in combination with sulfur-type vulcanization additives were studied using HNBR with 43 wt.% AN [35]. When primary or quaternary ammonium salts served as clay intercalants, the clay was intercalated in the uncured compounds. Moreover, the primary amine yielded higher interlayer spacing than the quaternary one (i.e., 3.85 nm compared to 3.53 nm). However, by sulfuric curing the organoclay bearing primary amine was confined to 3.10 nm and presented at the same time a second interlayer spacing at 1.70 nm (not connected with higher order reflections). On the other hand, the organoclay modified by the quaternary ammonium salt expanded further during vulcanization (3.87 instead of 3.53 nm) and XRD data suggested the presence of regularly intercalated structures. This behavior was traced to the interference of the primary amine with the sulfuric curatives during the creation of Zn complexes [45]. According to the related proposal the primary amine leaves the silicate surface and participates in the formation of the vulcanization intermediates in form of Zn complexes. Although this process may promote nanocomposite formation via interlayer cross-linking, the migration of the initial intercalant (whereby Zn2þ ions overtake the charge balancing role) to the matrix results in deintercalated organoclay structures. As shown in Figure 13.3, the TEM images depict the organoclay (i.e., layered silicate modified by ODA) confinement in HNBR after vulcanization. The observed interlayer opening in Figure 13.3a produced after melt blending of HNBR with the organoclay in an inner mixer was diminished after sulfur-type curing (Figure 13.3b). Such deintercalation phenomena were hindered by curing with peroxide wherein vulcanization involves thermal decomposition of peroxide generating at the same time free radicals on the rubber chains followed by their recombination [36]. Nevertheless, the vulcanization temperature and pressure values may cause the reordering of the initially disordered clay platelets in NBR/organoclay systems [46].

13.4 PROPERTIES 13.4.1 Mechanical (Dynamic–Mechanical) Addition of the appropriate type and amount of organoclay in an NBR mix enhances the mechanical performance of the respective vulcanizate. Such behavior has been

PROPERTIES

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FIGURE 13.3 TEM images of HNBR reinforced with layered silicate (FHT, SomasifÒ ) modified by octadecylamine (ODA) after (a) melt blending and (b) sulfur curing.

presented for a solution mixed vulcanizate (peroxide cured) of NBR bearing 50 wt.% AN content, for which the tensile strength and elongation at break values were saturated after 6 phr of organoclay loading at about 7 MPa and 800%, respectively [47]. The trend of property enhancement may be different as a function of the NBR polarity. For example, NBR with 34 wt.% AN content exhibited the highest tensile strength and strain at break values (i.e., 4.8 MPa and 920%, respectively) at

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RUBBER–CLAY NANOCOMPOSITES BASED ON NITRILE RUBBER

4 phr loading of solution compounded organoclay in the respective peroxide-cured vulcanizates [47]. However, the preparation method lends one more parameter in governing the mechanical performance of a rubber–clay nanocomposite [48]. Increase of the AN content from 26 to 35 and 42 wt.% in a sulfur-cured melt blended NBR/organoclay stock (10 phr filler amount) yielded tensile strength values of 3.7, 10.8, and 16.9 MPa, respectively. Similarly, the tear resistance was enhanced with the polarity increase, unlike the elongation at break values, which showed an opposite trend [27]. The mechanical property improvement is not always in concert with the basal spacing expansion of the organoclay. In a solution mixed NBR (32 wt.% AN content) with Mt modified by dimethyl dehydrogenated tallow quaternary ammonium salt, the tensile strength continuously increased up to 13.5 MPa as the organoclay reached 10 phr. At the same time the interlayer spacing reduced from 4.26 nm at 3 phr organoclay to 3.53 nm at 10 phr filler loading. The dispersed intercalated organoclay stacks, which were responsible for this tensile performance, yielded additionally a significant improvement in the tear resistance [17]. Effects of clay loading (up to 30 phr) were investigated by Wu et al. [21] in NBR–clay nanocomposites (AN content of the NBR: 24–26 wt.%) produced via the latex route. In spite of the marginal interlayer expansion of the clay, the dispersion of silicate stacks in the rubber matrix was able to endow to the vulcanizate a tensile strength of about 18 MPa. On the other hand, a tensile strength maximum value has been also reported at low clay concentration for NBR–clay nanocomposites prepared by the latex method. Hwang et al. [22] found that the tensile strength for NBR–Mt nanocomposites (AN content of the NBR: 31–33 wt.%) reaches a maximum (7.5 MPa) at about 5 phr of clay loading. In the same series of samples, a maximum tear strength was noticed for about 8 phr clay content. The clay dispersion, responsible for the strengthening of the related rubber–clay nanocomposites at room temperature during a tensile test, is expected to affect the performance of the nanocomposites in the whole temperature range. As detected by means of dynamic mechanical analysis (DMA) (Figure 13.4a), the NBR–clay nanocomposites exhibited higher storage moduli (E0 ) in the glassy state, as well as, in the rubbery one compared the neat matrix [23]. Note that these NBR vulcanizates with 34 wt.% AN content have been produced via the latex route. Increasing the clay loading in the respective vulcanizates resulted in a constant upgrade of the E0 . Moreover, 10 phr of clay content yielded a higher storage modulus in the whole temperature range than that of the vulcanizate filled with 50 phr of carbon black (CB). At the same time, incorporation of clay in NBR restricted the mobility of the matrix lowering the damping of the nanocomposites (Figure 13.4b). Although both CB at 50 phr and clay at 10 phr reduced the loss tangent (tan d) values at the glass transition temperature (Tg), only the latter shifted the Tg toward higher temperatures (Tg ¼  13.7 and  9.0
C for plain and clay reinforced NBR, respectively) [23]. This fact is related with the different interphase quality, created by CB and clay, respectively, within the NBR matrix. Note that the damping performance of NBR–organoclay nanocomposites depends on the AN content of the parent NBR. Sadhu and Bhowmick [49] presented for a nitrile rubber vulcanizate filled with 4 phr

PROPERTIES

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(a) NBR NBR-2M NBR-5M NBR-10M NBR-50CB

E' (MPa)

103

102

101

100 –50

0

50

100

150

Temperature (ºC) (b) 2.0 NBR NBR-2M NBR-5M NBR-10M NBR-50CB

1.6

Tan δ

1.2 0.8

0.4 0.0 –50

0

50

100

150

Temperature (ºC)

FIGURE 13.4 (a) Storage modulus (E0 ) and (b) loss tangent (tan d) versus temperature for NBR and its nanocomposites. Designations: 2, 5, and 10M represent 2, 5, and 10 phr organoclay contents, respectively, whereas 50CB designates the filling with 50 phr carbon black [23]. Copyright 2006, with permission from Springer Science þ Business Media.

organoclay that the tan d peak was shifted to lower temperatures for an NBR with 19 wt.% AN content, whereas Tg shift was noticed toward higher temperatures for an NBR with 50 wt.% AN content compared to each respective plain stock. Kim et al. [30] investigated the effect of various clay intercalants on the mechanical properties of NBR–organoclay nanocomposites. The storage modulus of the vulcanizate in the rubbery region was superior when ODA served as clay intercalant compared to the compounds with organoclays either with dodecylamine or octylamine modifications. Similarly, ODA intercalant delivered the best tensile performance (tensile strength 14 MPa and elongation at break 1500%) above 10 phr organoclay content. The mechanical performance of NBR–organoclay nanocomposites, cured by sulfur, was beneficially influenced by the excess amount of SA (i.e., 4 phr). This was ascribed to the presence of a fine clay dispersion in terms of

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extensive intercalated and exfoliated clay populations [32]. Liu et al. [50] argued that the mechanical performance of NBR–organoclay nanocomposites may further be improved via reactive mixing. For this purpose, NBR was compounded with organoclay in presence of resorcinol and hexamethylene tetramine complex. The reaction between the compounding components was made responsible for the shift in the Tg of the vulcanizates to higher temperatures, as detected by DMA [50]. Silane coupling agent was found to act synergistically on the mechanical performance of NBR–organoclay nanocomposites. Incorporation of 2–5 phr 3(mercaptopropyl)trimethoxysilane in a sulfur-cured NBR (29 wt.% AN content) yielded a tensile strength of about 13 MPa. Above this silane concentration, however, property deterioration was detected [51]. In order to enhance the mechanical behavior of NBR in the whole temperature range under demanding loading conditions HNBR [52] and HNBR–clay nanocomposites [53] are the right choices. For example, sulfur-cured HNBR–organoclay (10 phr) nanocomposites yielded a significant storage modulus increase, especially in the rubbery regime, compared to the plain matrix. At the same time, the mechanical damping (loss factor) of the nanocomposites was reduced, particularly for the organoclay bearing a quaternary intercalant with hydroxyl groups that was supposed to create hydrogen bonds with the AN groups of the rubber [35]. Apart from the effect of clay intercalant on the mechanical properties, the clay itself plays an important role. Mixing of HNBR with fluorohectorite (FHT) instead of Mt, both modified by ODA, was found to increase the modulus (stiffness) of the respective vulcanizates [54]. This is due to the significantly higher aspect ratio of FHT compared to Mt. It should be mentioned that the mechanical enhancement of HNBR–organoclay nanocomposites, compared to the neat matrix or conventional reinforced vulcanizates, is preserved even after aging at elevated temperatures in air, water, or oil [33]. The tensile strength improvement in HNBR–organoclay nanocomposites was accompanied by an increase of the strain at break values, which is very unusual for rubber–clay nanocomposites. This is, however, a well-described phenomenon in filled thermoplastic polymers. According to the related analogy, subcritical cracks induced in the vicinity of the platelets (via voiding, dewetting phenomena, chain slippage, etc.) enhance the stretchability (ductility) of the corresponding nanocomposite. Moreover, the zigzag route of a crack, as it propagates around the silicate platelets, assists further the energy dissipation mechanism. Thus, the nanocomposite may exhibit higher strain values than the neat matrix [35]. Videomonitoring the uniaxial tension test, Ahmadi et al. [55] detected macroscopic dilatation in NBR–organoclay nanocomposites. This supports that cavitation may be at work, in fact, in the vicinity of the clay platelets during loading as supposed above. Significant mechanical improvement has been also reported for LDH reinforced XNBR. LDH modified by decane sulfonate was melt-mixed with XNBR containing 7 wt.% carboxylic acid. Their vulcanizates showed an increase in the modulus values at all elongations when the organoclay content increased up to 10 phr [56]. Additionally, the mechanical property upgrade caused by the modified LDH was demonstrated by comparing the DMA spectra of the nanocomposite and the neat rubber matrix [39,56].

PROPERTIES

421

The mechanical performance of organoclay-reinforced NBR and XNBR has been modeled by Wu et al. [57], who introduced in the classical composite equations of Halpin–Tsai, modified Halpin–Tsai and Guth a modulus reduction factor (MRF). This factor took into consideration that a platelet-like filler has lower contribution to the Young’s modulus than a rod-like one. By fitting various experimental data the MRF value was estimated to be 0.66. It turned out that the modified Halpin–Tsai and Guth equations are the best description when this MRF term is considered. Following this approach, the modulus of HNBR– organoclay nanocomposites was modeled using the Guth equation with MRF extension [54]. 13.4.2 Friction/Wear Numerous applications of parts from nitrile rubbers, such as roll covers, outsoles, conveyor and timing belts, experience various modes of wear during service [3,58]. Therefore, NBR–clay nanocomposites should also be able to withstand similar demanding wear conditions. NBR reinforced with organoclay has been tested for abrasion. A cylindrical shaped NBR specimen was abraded on an abrasive sheet, which was placed on the surface of a rotating drum. Incorporation of 10 phr organoclay along with conventional CB filler increased the abrasion resistance of the plain matrix. However, the best performance was found when the organoclay loading increased up to 20 phr. Note that the same specimen showed also the lowest compression set [59]. The orientation of the intercalated clay platelets within the matrix is expected to affect the properties of the respective nanocomposites [34]. Thus, wear cannot be an exception in this respect. Calendering and hot pressing of a rubber/organoclay compound can actually result in an in-plane alignment of the clay platelets. Such a two-dimensional orthotropic character of a vulcanizate, wherein the clay layers are oriented parallel to the sliding direction, has been presented for HNBR–organoclay nanocomposites. It was found that this specific clay orientation acts disadvantageously on the sliding wear during a pin-on-plate (POP) type sliding wear test. This was reasoned by assuming a “can opening-type” mechanism. This mechanism is at work during the compression–extension strain cycle at the slider contact, and especially in the tensile zone at the rear part of the pin [60]. Nevertheless, in the respective HNBR/organoclay vulcanizates the coefficient of friction (COF) increased by almost 20% compared to the neat matrix. Apart from the dispersion degree of layered silicates in rubber–organoclay nanocomposites, the wear testing methods themselves affect the wear properties in large extent [61]. HNBR has been tested upon dry rolling and sliding using various setups as shown in Figure 13.5. As the slider moves over the rubber surface it creates wrinkles, which fold back while the indenter passes by. These “waves of detachment” (also called “Schallamach waves”) hamper the true sliding between surfaces creating a “saw tooth-like” profile on the abraded rubber [62]. Incorporation of 10 phr organoclay in HNBR decreased the width of the Schallamach pattern and the distance between the neighboring waves during an orbital rolling ball-on-plate

422

RUBBER–CLAY NANOCOMPOSITES BASED ON NITRILE RUBBER

FN

Counterpart (steel ball) Wear track

Orbital-RBOP

Sample (rubber) Rolling

FN

POP Counterpart (steel cylinder)

Counterpart (steel pin)

FN

ROP

Rotation Sample (rubber)

Wear track Sample (rubber) Sliding

FIGURE 13.5 Schemes of the set-ups of rolling (orbital-RBOP) and sliding wear (POP, ROP) tests [63]. Copyright 2010, with permission from Springer Science þ Business Media.

(orbital-RBOP) test. At the same time, the coefficient of friction was enhanced whereas the specific wear rate was reduced [63]. As shown in Table 13.1, under sliding conditions the COF was less affected in POP while reduced during roller-onplate (ROP) configurations (Figure 13.5) for the nanocomposites compared to the plain HNBR matrix. On the other hand, the wear rate of the rubber–organoclay nanocomposite was reduced under POP and ROP compared to the unfilled reference HNBR. This was attributed to changes in the failure mechanisms as disclosed in Ref. [63]. TABLE 13.1 Hardness and Wear Properties of HNBR and HNBR/Organoclay Compounds Property/Test Hardness Rolling (orbital-RBOP) COF Specific wear Sliding (POP) COF Specific wear Sliding (ROP) COF Specific wear Source: After Ref. [63].

Unit Shore A

HNBR Reference

HNBR þ 10 phr Organoclay

42

52

mm /(N m)

4.2  10  2 3.0  10  4

5.4  10  2 1.3  10  4

mm3/(N m)

1.14 1.1  10  1

1.24 3.7  10  2

mm3/(N m)

3.00 3.8  10  2

1.89 7.8  10  3

3

PROPERTIES

423

13.4.3 Barrier The most popular property enhancement in case of rubber–layered clay nanocomposites is considered to be the gas barrier performance [11]. Hence, lot of research has also been conducted on this topic regarding nitrile rubbers. NBR–clay nanocomposites prepared through the solution method reduced significantly the water and methanol vapor permeability. Noteworthy is the 80% water vapor permeability decrease just by adding 1 phr organoclay. On the other hand, the nanocomposites started to act for methanol vapor barrier only after 3 phr organoclay loading [17]. The platelet network created in the NBR matrix acts as sufficient gas barrier also for vulcanizates prepared via the latex route. Wu et al. [24] using 20 phr of Mt found a reduction of about 50% in the nitrogen permeation of the nanocomposites compared to the neat matrix. Barrier improvement owing to the silicate platelets becomes prominent when nanocomposites are compared with microcomposites containing conventional isotropic fillers. The latter have limited barrier action even at high filler contents [21]. However, other sealing applications might require a combined action of spherical and platelet-like fillers. This has been proved for melt blended NBR, containing both CB and organoclay, in respect to the nitrogen permeability [59]. The clay-induced barrier improvement may be enhanced by using suitable coupling agents. Kim et al [41] varied the amount of 3-(mercaptopropyl)trimethoxysilane in a mix of NBR with 5 phr organoclay. It was found that the initially low water vapor permeability of the nanocomposite was further reduced by increasing the concentration of the silane coupling agent. This is likely a combined effect of optimized clay dispersion and interphase quality. On the other hand, by varying the amount of AN in NBR the barrier performance has been changed, as well. As shown in Figure 13.6, NBR with increasing AN content becomes less susceptible to the permeation of nitrogen molecules. However, incorporation of 10 phr organoclay in the compounds

35 Pure 26NBR

P [10–17 m2 (Pa.s)–1]

30

26NBR/OCNs

25 20 15

Pure 35NBR 35NBR/OCNs

10

Pure 42NBR 42NBR/OCNs

5 0

FIGURE 13.6 Nitrogen permeability (P) of NBRs with different AN contents and related NBR–organoclay nanocomposites containing 10 phr organoclay (OCN) [27]. Copyright 2009, with permission from John Wiley and Sons.

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RUBBER–CLAY NANOCOMPOSITES BASED ON NITRILE RUBBER

caused a comparable barrier enhancement irrespective to the AN content [27]. This may be argued by assuming that the tortuous path created by the platelet networks in the NBR matrices of varying AN contents, is similar to one another. Variation of the tortuosity of the penetration path owing to the impermeable silicate layers affects the barrier property of the respective rubber nanocomposites. HNBR reinforced with 10 phr Mt modified by methyl-tallow-bis(2-hydroxyethyl) quaternary ammonium salt exhibited intercalated structures of high basal spacing. Moreover, the formation of a strong interphase was suspected due to hydrogen bonding in the related peroxide-cured vulcanizate [36]. The above phenomena, that is, intercalation with high basal spacing and formation of strong interphase, may further reduce the permeability of the related rubber/organoclay compounds. Further gas barrier improvement may be derived by increasing the aspect ratio of the silicate sheets. Using high aspect ratio FHT instead of low aspect ratio Mt, both modified by the same ODA intercalant, the oxygen transmission rate could be further reduced [54]. Note that both silicates had similar interlayer spacing, while the Mt possessed significantly lower aspect ratio than the FHT. As expected, unmodified FHT at the same filler loading (10 phr) has limited gas barrier improvement compared to the plain matrix due to the existence of FHT agglomerates (i.e., microcomposite structure). The gas barrier performance of HNBR–organoclay nanocomposites has been successfully described by using either the Nielsens’ or the Bharadwajs’ equations [54]. The dispersion of organoclay in NBR affects additionally the swelling performance of the respective vulcanizate. NBR reinforced with Mt modified by ODA was less swollen in MEK compared to the compound with the same clay fraction, in which the Mt was modified by octylamine or dodecylamine [29]. Similarly, the dispersion of the organoclay platelets in HNBR affects the solvent resistance of the vulcanizates [64]. 13.4.4 Fire Resistance Applications of NBR products for gas and fuel piping or under-hood automotive parts demand enhanced fire resistance. NBR reinforced with 20 phr clay, which was prepared via the latex route, was investigated for its flame retardance [20]. The respective sulfur-cured vulcanizates were tested in cone calorimeter using an incident heat of 50 kW/m2. As shown in Figure 13.7, the NBR–clay nanocomposite exhibited the lowest peak of heat release rate compared to the neat matrix or to conventional microcomposites. Note that to produce NBR–clay microcomposite the clay was incorporated in 20 phr directly on an open mill along with the curatives. The differences among the investigated vulcanizates regarding the ignition time were rather marginal, while the peak of the heat release rate was shifted toward longer time for the nanocomposite compared to the matrix, silica-filled nanocomposite and clayfilled microcomposite [21]. Nevertheless, a combination of fillers such as clay and carbon nanotubes (CNT) has been quoted to have a synergistic effect on the thermal degradation of HNBR. This was attributed to the higher char production in the ternary hybrid (HNBR–clay–CNT) compared to the binary nanocomposites [65].

OUTLOOK

425

2400 Pure NBR NBR–clay nanocomposites NBR–clay microcomposites NBR–silica composites

Heat release rate (KW/m2)

2100 1800 1500 1200 900 600 300 0

0

50

100

150

200

250

300

Time (S)

FIGURE 13.7 Comparison of the heat release rate versus time plots for pure NBR and its composites [20]. Copyright 2008, with permission from John Wiley and Sons.)

13.4.5 Others The recent development with NBR–clay nanocomposites focuses mostly on the enhancement of “traditional” properties (i.e., mechanical, friction, barrier), in respect of which the corresponding conventional microcomposites have certain limitations. On the other hand, incorporation of layered clay in (H)NBRs was found to improve other properties, as well (e.g., reduction of the shear viscosity and the die-swell of the clay-reinforced compound compared to the neat matrix). Moreover, the gradual increase of the organoclay loading reduced further the die-swell of the corresponding nanocomposites [18]. XNBR reinforced with 10 phr organoclay has been found to enhance the relaxation strength of the high temperature relaxation process compared to the plain matrix [66]. This characteristic, likely due to the mobile ions on the clay particle surface or inside the filler, leads to Maxwell–Wagner polarization effects, which were not observable in the respective silica filled compounds [38]. In several cases, clay has been exploited as compatibilizer in immiscible rubbers. In such an example NBR was mixed with styrene–butadiene rubber (SBR) by simultaneous incorporation of Na-Mt in the compound. The clay was identified to be located at the interface between the immiscible blend components [67].

13.5 OUTLOOK Polar butadiene-type rubbers (NBR, HNBR, XNBR, and so on) are excellent model materials for organoclay modification owing to their polarity. Research works on clay

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RUBBER–CLAY NANOCOMPOSITES BASED ON NITRILE RUBBER

containing nanocomposites will next concentrate on some of the long-term properties, such as fatigue performance, resistance to ozone cracking, and aging under different conditions. The results achieved so far allow us to conclude that the organoclay incorporation will be favored for improvement of the barrier and fire resistance properties. Considerable effort will be overtaken to meet the above target by combined use of organoclays and novel (such as carbon nanotubes) and especially traditional fillers.

ACKNOWLEDGMENT This work is connected to the scientific program of the “Development of qualityoriented and harmonized RþD þI strategy and functional model at BME” project. This project is supported by the New Hungary Development Plan (Project ID: ´ MOP-4.2.1/B-09/1/KMR-2010-0002). TA

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CHAPTER 14

RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS C. PHILIPPE MAGILL DANA K. ADKINSON RALF I. SCHENKEL

14.1 INTRODUCTION 14.1.1 Butyl Rubber: Key Properties and Applications Butyl rubber is a member of the family of isobutylene-based polymers that includes polyisobutylene (PIB), copolymers of isobutylene and isoprene (butyl rubber, IIR), halogenated butyl rubber (XIIR, where X ¼ Cl or Br) and brominated copolymers of isobutylene and para-methylstyrene (BIMSM) [1,2]. Their chemical structures are shown in Figure 14.1. Terpolymers of isobutylene, isoprene, and either divinylbenzene or para-methylstyrene are also known, as are star-branched butyl polymers (SBB) containing 5–20% of a highly branched fraction. All of these polymers will be included within the context of the ensuing chapter, and for the sake of simplicity, the term “butyl rubber” will often be used generically in the text to encompass the entire family of isobutylene-based polymers. A unique combination of properties makes butyl rubber commercially important: . . . .

outstanding impermeability to air and moisture; very good damping properties (low resilience); good low temperature flexibility due to its low glass transition temperature (Tg  70 C); good resistance to heat, oxidative and chemical degradation due to its highly saturated polymer backbone.

Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

431

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RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

(a) Polyisobutylene (PB)

97.5–99.2 mol%

0.8–2.5 mol%

(b) Butyl rubber, poly(isobutylene-co-isoprene) (IIR) X

(c) Halogenated butyl rubber (XIR), X = Cl, Br

(<5 mol%)

(<2 mol%)

Br

(d) Brominated poly(isobutylene-co-paramethylstyrene) (BIMSM)

FIGURE 14.1 Chemical structures of (a) polyisobutylene, (b) butyl rubber, (c) halogenated butyl rubber, and (d) brominated poly(isobutylene-co-para-methylstyrene).

This combination of properties is unique among elastomers. Other polymers may exhibit low permeability and low resilience, such as nitrile rubber, but have relatively high Tg’s and have poor low temperature flexibility, while polymers with low Tg’s, such as polybutadiene and polyisoprene exhibit very high permeability to air. The impermeability of butyl rubber is largely attributable to the close packing of the isobutylene chain segments of which it is primarily composed [3]. Two methyl side groups on every other carbon atom in the backbone cause steric hindrance. The chains therefore move relatively slowly so that they block the passage of gas molecules rather than moving aside and letting them pass. The bulky methyl groups also prevent the chains from taking up a regular configuration and crystallizing, while at the same time they have only small dipole moments, which do not give rise to strong interchain attractions. Because of this structure, butyl rubber has good low temperature flexibility and high damping characteristics. It is the impermeability of butyl rubber to air that mainly accounts for its practical utility. Greater than 80% of all butyl and halobutyl rubber goes into tire innerliners and innertubes, since it is the material of choice for maintaining tire inflation pressure. Other applications requiring impermeability for which butyl rubber is used include tire curing bladders, pharmaceutical stoppers, chemical tank linings, air conditioning hose, protective suits, boots, and gloves. The chemical and heat resistant properties of butyl rubber are an additional benefit in many of these

INTRODUCTION

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applications. With respect to its vibration damping characteristics, butyl rubber finds use in applications such as automotive body mounts, antivibration pads, and pneumatic air springs. In the tire industry, there is a continuing drive toward improving the air retention capability of tires, since a properly inflated tire is essential to achieving optimum tire performance. Factors such as rolling resistance that have an impact on vehicle fuel economy, and heat build-up in the tire which impacts tire wear and durability are dependent on correct tire inflation pressure. At the same time, there is a desire to reduce tire weight and material costs. It is not surprising therefore that butyl and halobutyl rubber nanocomposites have become a subject of great interest to the tire industry since they hold the promise of greatly improved air barrier properties. 14.1.2 Butyl Rubber–Clay Nanocomposites Clays have long been used in butyl rubber compounds as low cost fillers. Their principle benefits have been in reducing the overall compound cost, while providing low-to-moderate reinforcement and improved processing for extrusion and calendering [4]. Other compound properties may suffer, however, such as tear strength and adhesion or tack, particularly when high fillers loadings are used. These deficiencies may be overcome with the use of nanoclays due to the much greater level of reinforcement that is possible at relatively low filler loading levels. Butyl rubber–clay nanocomposites may be grouped into three different categories depending on the degree of clay dispersion within the compound (Figure 14.2). .

Conventional Microcomposites (Figure 14.2a). The polymer and clay are phase separated, due to poor compatibility of the two phases. Polymer chains are not

Polymer

Layered silicate

(a) Phase separated (microcomposite)

(c)

(b) Intercalated (nanocomposite)

Exfoliated (nanocomposite)

FIGURE 14.2 Principal morphologies of polymer–clay composites. Reproduced from Ref. [5] with permission from Elsevier.

434

.

.

RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

able to wet the surface of the clay due to a large difference in surface energies, and consequently no polymer will penetrate the clay interlayer spaces. Very little bonding interaction exists between the clay and the polymer resulting in little mechanical reinforcement. Butyl rubber composites with unmodified inorganic clays mostly fall into this category. Intercalated Nanocomposites (Figure 14.2b). Polymer chains are able to penetrate the clay layers, increasing the spacing between the layers. Bonding interaction between the polymer and clay results in some reinforcing effects, although not to an optimum level. Organically modified clays that are more compatible with butyl rubber must generally be used in order for polymer intercalation to occur. Exfoliated Nanocomposites (Figure 14.2c). An exfoliated nanocomposite is characterized by the clay layers being individually separated and randomly dispersed within the polymer matrix. For this morphology to be achieved, the polymer chains must first intercalate and then separate the clay layers. In addition to very good compatibility between the clay and the rubber, high mixing shear is normally required to achieve clay exfoliation. This morphology results in a high level of polymer reinforcement and excellent barrier properties.

There is some dispute in the literature as to whether a fully exfoliated clay morphology is actually attained in practice [6,7]. The true state is most likely a mixture of intercalated and exfoliated morphologies. As with other nonpolar olefinic polymers, obtaining the desired level of clay exfoliation in butyl nanocomposites is a significant challenge due to the hydrophobic nature of the polymer. Consequently, much of the research into butyl rubber–clay nanocomposites deals with methods to compatibilize the clay and the polymer, as will be discussed in Section 14.3. 25 Papers Patents

20

15

10

5

FIGURE 14.3 Chronology nanocomposites.

of

publications

and

patents

related

2010

2009

2008

2007

2006

2005

2004

2003

2002

2001

2000

1999

1998

1997

1996

1995

0

to

butyl–clay

TYPES OF CLAYS USEFUL IN BUTYL RUBBER–CLAY NANOCOMPOSITES

435

The total body of published literature devoted to butyl rubber–clay nanocomposites is relatively small, but has been growing over the past 15 years (Figure 14.3). The patent literature surpasses the academic literature by a wide margin, which can be seen as a reflection of the commercial interest surrounding this technology. Much of the literature is focused on improving the barrier properties of butyl rubber, particularly in tire-related applications, as will be discussed in Section 14.5.

14.2 TYPES OF CLAYS USEFUL IN BUTYL RUBBER–CLAY NANOCOMPOSITES A review of the literature shows that a number of different types of clays have been explored with butyl rubbers. As expected, the majority of studies have been done with montmorillonite clays since they are the most widely known. Other clays that have been investigated include vermiculite, sepiolite, mica, synthetic hydrotalcite clays, and high aspect ratio talc fillers. 14.2.1 Montmorillonite Clays Montmorillonite clays fall within the phyllosilicate family of layered silicate clays (phyllo meaning sheet-like). The crystal structure consists of layers made up of two silica tetrahedra edge-shared with an octahedral sheet of either alumina or magnesia in a 2:1 fashion. The sheets are typically 100–200 nm in length, and 1 nm in thickness, resulting in aspect ratios of between 100:1 and 200:1. Montmorillonite clays are commercially available either in their natural unmodified form (sodium montmorillonite, Na-Mt), or in organically modified variants (abbreviated here as OMt). Organic modification of montmorillonite clays renders the clays more organophilic, and is generally considered to be a necessary step in order to make them compatible with nonpolar polymers such as butyl rubber [6,8]. 14.2.2 Hydrotalcite Clays Hydrotalcite clays, also referred to as layered double hydroxides (LDH) belong to a class of anionic clay minerals in which anions such as OH  , Cl  , CO32  , NO3  , and so on occupy the spaces between positively charged sheets of Mg/Al hydroxide. The chemical formula of the most common form of hydrotalcite is given as Mg6Al2(OH)16CO34H2O. The minerals can be both naturally mined or synthetically made. The synthetic materials are said to be free of metallic impurities, which can cause polymer degradation and are claimed to have superior high temperature stability compared to naturally obtained clays [9]. The inorganic anions are readily exchanged with bulky organic anions in order to enlarge the interlayer spacing and facilitate the intercalation of polymer chains. Hydrotalcites are characterized by having a high aspect ratio of between 200:1 and 400:1, which makes them good candidates for enhancing polymer barrier properties.

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RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

Their metal hydroxide layers undergo endothermic decomposition on heating, which can improve the flame retardancy of polymer nanocomposites [10]. Bhowmick and coworkers carried out studies to evaluate the interaction of LDH fillers with polar polymers such as carboxylated nitrile rubber (XNBR) [10], polyurethane (PU) [11], and ethylene vinyl acetate (EVA) [12–14] versus nonpolar polymers such as EPDM [10,15]. The hydrotalcite clay was modified with long chain anion surfactants, such as sodium decanesulfonate (C10H21SO3Na) or sodium dodecylsulfate (C12H25SO3Na). They observed strong interactions between the clay and the polar polymers but very poor affinity for the nonpolar polymers, such that the LDH clays were barely wetted. Sakazaki and colleagues (Tokai Rubber Industries) describe a butyl rubber–LDH nanocomposite formulation useful in a low permeability gasohol fuel filler hose application [16]. The authors report that when an outer hose layer is prepared comprising a blend of butyl and bromobutyl rubber with 5 parts per hundred rubber (phr) of a hydrated hydrotalcite clay (DA-500 from Daiso Co., Ltd.) and 2 phr of 1,8diazabicyclo(5,4,0)undecene-7 salt (DBU salt), there is an improvement in interlayer adhesion between the butyl rubber outer layer and the innertube polymer (NBR–PVC, polyamide, or fluoropolymer). This is said to be due to a synergetic effect of the acid accepting behavior of the hydrated hydrotalcite compound and the vulcanization accelerating effect of the DBU salt. Alcohol permeation of the hose was reduced by as much as a factor of 20 when the butyl rubber composition was used in place of NBR–PVC. Winters et al. (Akzo Nobel) describe the modification of hydrotalcite with a charge balancing organic anion [17]. The organic anion used is a long chain fatty acid (palmitic acid, C16), which when intercalated between the clay layers, results in an  interlayer spacing increase from 8 to 29 A. If an unsaturated acid is used, the pendant double bonds can participate in sulfur cross-linking reactions with the polymer. The composites exhibit a 40–50% improvement in tensile strength, elongation, and tear strength compared to the base polymer compound. Alkoxy silane coupling agents such as bis(3-triethoxysilylpropyl) tetrasulfide (Si69, Degussa) can also be used to increase the bonding interaction between the clay and the polymer. Ebner and Hallock (Cryovac Inc.) describe the use of a hydrotalcite clay modified with a transition metal salt that resulted in a composition having oxygen scavenging properties [18]. The composition is useful in food or pharmaceutical packaging (films or sealing gaskets). The modified hydrotalcite clay is dispersible in a polymeric carrier that may include butyl rubber. Overall, the amount of published data related to butyl–hydrotalcite nanocomposites is too small to draw any clear conclusions. But given its high aspect ratio relative to other clays, and its ability to be modified with a wide array of organic anions, hydrotalcites in combination with butyl rubbers are worthy of further exploration. 14.2.3 High Aspect Ratio Talc Fillers Talc is a low cost nonreinforcing filler that is widely used in the rubber industry. It is often added to carbon black filled compounds to improve processability [19]. Talc

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is a hydrated magnesium silicate of formula Mg3Si4O10(OH)2 that is neutral in charge having no interlayer cations or anions. Rio Tinto Minerals has developed a process to produce a laminar talc with a relatively high aspect ratio of 20:1 and a surface area of between 15 and 25 m2/g [20]. In their promotional literature, Rio Tinto claims that incorporation of 30 phr of high aspect ratio talc in a bromobutyl innerliner formulation results in a 40% reduction in air permeability [21,22]. This is to some extent confirmed by Krueger in a recent patent filing in which 30 phr of high aspect ratio talc was substituted for carbon black in a tire innerliner formulation [23]. The air permeability of the innerliner compound was reduced by 30–38%. Flex fatigue properties were also seen to improve. No results are given with regard to adhesion, which in tire construction is a critical property for lamination of the butyl innerliner to the tire carcass. Adhesion is likely to suffer when such high levels of talc are present in the formulation. Resendes and Adkinson at LANXESS Inc. investigated several talcs with aspect ratios between 1:5 and 1:7 in bromobutyl formulations [24]. A comparison was made between a standard bromobutyl grade and novel grades containing higher levels of unsaturation, and modified with either an organic phosphine or amine to form high unsaturation bromobutyl ionomers. When 60 phr of talc was used in place of carbon black in a bromobutyl innerliner formulation, impermeability improved by approximately 25%, but modulus and tensile strength suffered, which is not surprising given the high talc loading. When high unsaturation bromobutyl ionomers were used in combination with 60 phr of talc, permeability decreased a further 20%, and modulus also increased, which is indicative of better compatibility between the talc filler and the ionic functionalized polymer, leading to better filler dispersion. The best results were obtained when a combination of talc (40 phr) and carbon black (20 phr) was used with a high unsaturation bromobutyl ionomer, in which case permeability decreased by as much as 60% compared to the standard bromobutyl formulation. Butyl rubber ionomers are discussed further in Section 14.3.2. 14.2.4 Other Clays Mica comprises a family of platy aluminosilicates of which muscovite with chemical formula (KF)2(Al2O3)3(SiO2)6(H2O) is the most common form. Like montmorillonite, they are 2:1 layered clay with alternating tetrahedral silicate layers and octahedral layers of magnesia and alumina. Micas are widely used fillers in the plastics industry and in the electronics industry for their electrical insulating properties. In several recent patents, Fudemoto and colleagues at Bridgestone Corp. give examples of organically treated mica being used in nanocomposites with bromobutyl rubber and BIMSM polymers [25–28]. Imidazolium surfactants were used to enhance the compatibility of the mica with the butyl polymers. The air permeability of the rubber compounds were reportedly reduced by up to 50% when 30 phr of imidazolium treated mica is incorporated. Vermiculite is a form of mica in which the potassium ions are replaced by magnesium and iron. The platelets can have an aspect ratio as high as 10,000:1 and they readily swell and disperse in water. This led to the selection of vermiculite as a

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RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

nanofiller for use in the butyl latex technology developed by InMat Inc. for barrier coating applications [29]. This technology will be discussed further in Sections 14.4.3 and 14.5.1. Sepiolite is a clay with a fibrous or needle-like structure which means that it does not strictly fit into the descriptive category of platy clays, although it is still considered a member of the phyllosilicate class. It is a hydrated magnesium silicate of general formula Mg4Si6O15(OH)26H2O and like the other phyllosilicate clays has an alternating 2:1 structure composed of two tetrahedral silica sheets and a central octahedral sheet of magnesia. The silica sheets tend to form Si–O–Si bonds along a single axis rather than a plane, resulting in a needle-like form. Nevertheless, sepiolite has a rather large surface area of 200–300 m2/g due to a very fine structure, which contains intricate channels and micropores. This lends the material to be useful in imparting reinforcing properties. Bhowmick and colleagues investigated the effect of sepiolite on the adhesive strength of brominated isobutylene-co-paramethylstyrene (BIMSM) [30]. With 2 and 8 phr of sepiolite clay in the formulation, the adhesive tack strength of the rubber increased by 140 and 300%, respectively compared to the base polymer. The elastic modulus was also enhanced. The authors posit that due to the fine fibrous nature of the clay particles, the polymer molecules are able to reach across the particles and maintain a well-entangled network, while at the same time benefiting from the reinforcing effects of the filler.

14.3 COMPATIBILIZER SYSTEMS FOR BUTYL RUBBER–CLAY NANOCOMPOSITES Filler dispersion is a common problem in polymer compounding due to the interparticular forces that keep the filler particles from deagglomerating [31]. These forces of attraction become orders of magnitude greater in the case of nanofillers because of their very large surface area [32]. Natural clays are hydrophilic and are generally considered to be incompatible with nonpolar polymers such as butyl rubber. The preferred starting point is therefore to use organically modified clays in which some of the alkali metal ions are exchanged with organic ions, most commonly organic amines or phosphines. This has the combined effect of rendering the clay surface more organophilic, and enlarging the interlayer spacing to facilitate polymer intercalation. However, this may still not be enough to achieve the desired level of clay exfoliation and therefore other compatibilization techniques must be applied. The principle approaches that have been used with butyl rubber are 1. The use of surfactants or swelling agents to further separate the clay layers. 2. Modification of the butyl polymer chains by the addition of polar groups which have an affinity or can bond with the clay surface. Examples include butyl rubber ionomers, and maleic anhydride modified butyl rubber. 3. Introduction of a second compatibilizing polymer. This is typically a lower molecular weight polymer species with polar or reactive groups, which can

COMPATIBILIZER SYSTEMS FOR BUTYL RUBBER–CLAY NANOCOMPOSITES

439

readily intercalate within the clay layers while at the same time being compatible with the butyl rubber matrix. Table 14.1 gives a summary of the different compatibilizer systems that have been used in the preparation of butyl rubber–clay nanocomposites, which will be discussed in more detail. 14.3.1 Surfactants and Swelling Agents Organically modified montmorillonite clays (OMt) that have been ion exchanged with quaternary amines are commercially available from the major clay suppliers. Clays modified with quaternary amine surfactants having long alkyl chains (C18) are recommended when dealing with nonpolar polymers such as butyl rubber. This has been confirmed in a recent study by Samadi and Kashani [60] in which they evaluated a series of different organic amine clay modifiers and found that those with bulky hydrogenated tallow and benzyl groups on the modifying amines worked the best. Mang and Han have documented a similar trend [80]. Liang et al. report that butyl alcohol, oleic acid, and castor oil are all useful as swelling agents for OMt clays [78]. In their experiments, a concentrated dispersion was first made of the OMt clay in the swelling agent, then the dispersion was introduced into either a melt or solution blend with butyl rubber. The nanocomposites prepared with the swollen clay exhibited superior tensile strength and gas barrier properties compared to nanocomposites in which the OMt clay was melt compounded without the preswelling step. Li and colleagues report similar findings using stearic acid as the swelling agent [79].Upon addition of stearic acid to OMt, the interlayer spacing increased from 26 to 40 A. Stearic acid is a common ingredient in sulfur-cured butyl rubber formulations and is therefore convenient to use as a clay-swelling agent. 14.3.2 Butyl Rubber Ionomers Ionomers are ionic polymers that contain a small amount (typically less than 10%) of ionic functionality within the polymer backbone. The introduction of these ionic groups offers the potential for good interaction with polar clay surfaces. Ionomers of butyl rubber have been known for some time and they can be readily prepared via melt or solution processes [81–83]. Figure 14.4 shows the reaction of organic amines and phosphines with bromobutyl rubber and BIMSM to give their corresponding ionomers. Parent and coworkers reported the synthesis and characterization of bromobutyl rubber ionomers by reaction with either triphenylphosphine or dimethyloctylamine in a melt process [84]. The phosphine ionomer (IIR–PPh3Br) was found to be more stable than the amine analogue, which had a tendency to eliminate ammonium bromide and form conjugated diene butyl. In a subsequent paper, Parent et al. investigated the interaction of the phosphine ionomer with OMt clay [33]. The ionomer nanocomposite was prepared in an internal batch mixer and showed evidence of good clay exfoliation by wide-angle X-ray scattering (WAXS) and

440

Maleic and succinic anhydride modified polymers (PIB, IIR, IR, BR, BIMSM) Dicyclopentadiene resin Terpene hydrocarbon resin Phenolic resin Low mol. wt. polymers (PIB, IIR, BR, BIMSM, polybutene) Aminated liquid PIB Epoxidized styrene–butadiene–styrene Vinylbenzyl chloride SBR Imidazolium surfactants

IIR, BIIR, BIMSM, SBB

IIR IIR IIR, BIIR

IIR

IIR BIIR, BIMSM

Long chain fatty acid and/or silane coupling agent Surfactant wetting agents Butyl alcohol Stearic acid

Phosphonium ionomer of BIIR BIMSM amine ionomer

BIIR BIMSM

BIMSM BIIR, BIMSM IIR, BIMSM IIR, BIIR, BIMSM, SBB BIIR, BIMSM, SBB BIMSM

Compatibilizer

Polymer

Melt (internal mixer) Solution, melt (internal mixer)

OMt þ dimethylstearyl-amine Mica modified with long-chain imidazoles Hydrotalcite Vermiculite OMt OMt

Melt (internal mixer) Melt (internal mixer)

Na-Mt, OMt OMt

Latex Melt (two-roll mill) Melt (two-roll mill)

Solution

Melt (internal mixer) Melt (internal mixer) Solution, melt (internal mixer) Solution, melt (internal mixer)

Melt (internal mixer, extruder) Melt (internal mixer), solution/ emulsion Melt (internal mixer, two-roll mill), solution/emulsion

Processing Method

Na-Mt, OMt OMt Na-Mt, OMt Na-Mt, OMt

Na-Mt, OMt

OMt Na-Mt, OMt

Nanoclay

TABLE 14.1 Compatibilizers Used with Butyl Rubber–Clay Nanocomposites

[29,73–77] [78] [79]

[17]

[72] [27,28]

[69,70] [71]

[55,56] [57,58] [59,60] [61–68]

[48–54]

[24,33–35] [36–47]

References

COMPATIBILIZER SYSTEMS FOR BUTYL RUBBER–CLAY NANOCOMPOSITES

441

Br

AR3 (A=N or P)

Δ

AR3+ Br– (a) Bromobutyl ionomer formation

NR3 Δ AR3+ Br–

Br (b) BIMSM ionomer formation

FIGURE 14.4 Ionomers of (a) bromobutyl rubber and (b) BIMSM.

transmission electron microscopy (TEM). In comparison, when bromobutyl rubber was blended with OMt clay in the internal mixer, the clay did not exfoliate as indicated by distinct peaks in the WAXS profile (Figure 14.5). Stress–strain analysis showed a significant increase in polymer reinforcement in the case of the

Arbitary units

IIR–PPh3Br + 15 wt. % NR4+–MM

Arbitary units

BIIR + 15 wt. % NR4+–MM

1

3

5 2θ (º)

7

9

FIGURE 14.5 WAXS profiles of IIR–PPh3Br and BIIR nanocomposites of NR4 þ –Mt. Reproduced from Ref. [33] with permission from Elsevier.

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RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

bromobutyl–ionomer nanocomposite, confirming a strong interaction between the ionomer and the clay, but no reinforcing effect was seen with the nonionomeric bromobutyl nanocomposite. Surprisingly, there was no improvement seen in air impermeability with the ionomer nanocomposite. In a related patent, Parent et al. prepared bromobutyl–montmorillonite nanocomposites in an internal mixer and evaluated the effect of adding PPh3 to form IIR–PPh3Br in situ [34]. The resulting ionomer nanocomposite exhibited twice the tensile strength of the nonionomeric BIIR nanocomposite, indicating significant reinforcement. Resendes and Adkinson used a similar bromobutyl phosphine ionomer in combination with high aspect ratio talc fillers and measured up to 60% improvement in air impermeability in a tire innerliner formulation [24]. Ammonium ionomers of BIMSM (Figure 14.4b) have been prepared via solution and melt processes and have been extensively investigated as a platform for BIMSM–clay nanocomposites [36–47]. Tsou and Measmer carried out a series of experiments comparing bromobutyl rubber, BIMSM, and BIMSM ammonium ionomers in melt mixes with varying amounts of OMt clay [37]. From Gusev– Lusti equations and permeability data, the authors calculated projected aspect ratios of the clay dispersed in the polymers to determine the extent of clay exfoliation. From this they concluded that BIMSM is able to disperse the OMt–clay somewhat more effectively than BIIR, and ammonium functionalized BIMSM ionomers gave the best results overall. In related patent filings, Dias and colleagues at ExxonMobil Chemical Co. prepared BIMSM ionomers in solution by the addition of long chain alkyl amines (C18 tallow alkyl amines) [40,41]. The ionomers when mixed with montmorillonite clays in an internal mixer yielded nanocomposites with superior air permeability properties. Fudemoto and Wang at Bridgestone Corp. similarly prepared BIMSM ionomers in an internal mixer and were able to exfoliate amine-modified mica clays to a greater extent than with untreated BIMSM, as evidenced by X-ray diffraction (XRD) WAXS and SAXS (small angle X-ray scattering) [45]. It has also been shown that BIMSM ionomers can be made in situ in solution at the same time that the clay is introduced [42]. Weng and colleagues disclose a continuous process by which a BIMSM ionomer is formed in a hydrocarbon solution followed by the addition of either an aqueous slurry of Na-Mt or a hydrocarbon dispersion of OMt, which upon emulsification and drying gives a well exfoliated clay nanocomposite with very good air barrier properties [43,44]. Introduction of ionic groups into a polymer can result in an increase in the polymer viscosity due to ionic interactions between the polymer chains. This attraction is sufficient for the polymer to form ionic aggregates, which can act like a weakly crosslinked network structure [85]. Depending on the strength of the electrostatic attractions, the aggregates can dissociate to some extent upon heating (Figure 14.6). However, the increase in polymer viscosity can pose challenges in polymer processing, such as in mixing and calendering during tire innerliner manufacturing. Bergman et al. show that this problem can be overcome if the number of ionic sites is kept to a minimum and if they contain bulky alkyl groups, which can inhibit the ionic interactions [47].

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Δ

FIGURE 14.6 Thermally reversible aggregation of ionomeric clusters.

14.3.3 Maleic Anhydride-Grafted Polymers Grafting of maleic or succinic anhydride groups is a common method to introduce reactive functionality into olefinic polymers. The grafting reactions are normally done in a melt process such as using a twin-screw extruder, with the aid of a free radical catalyst such as an organic peroxide. A side reaction can be polymer degradation due to radical chain scission, which can be significant in the case of butyl polymers since the tertiary carbon atoms of the isobutylene groups are particularly susceptible to cleavage. Kato et al. report on the preparation and use of maleic anhydride-grafted butyl (MAH-g-IIR) in nanocomposites with montmorillonite clays [48]. Low molecular weight MAH-g-IIR was prepared in solution by reacting IIR and MAH in the presence of a peroxide initiator. The low molecular weight MAH-g-IIR polymer was then melt compounded with varying levels of OMt, as well as Na-Mt and carbon black filled compounds for comparison. The nanocomposites showed evidence of good clay dispersion by XRD and TEM analysis. The MAH-g-IIR organoclay nanocomposites showed superior tensile strength, modulus, and gas barrier properties compared to the Na-Mt and carbon black filled compounds. Osman prepared a MAH-g-BIIR nanocomposite and investigated its adhesive properties [54]. The MAH grafting reaction was carried out in solution using a novel metal alkanoate catalyst, which resulted in very little molecular weight breakdown. MAH grafting could similarly be done in the presence of clay. The adhesive strength of the maleated bromobutyl and its nanocomposite were measured in conjunction with various substrates, as discussed in Section 14.5.4. Maruyama et al. at Yokohama Rubber Co. used a low molecular weight liquid maleated polyisobutylene (MAH-gPIB) as a clay compatibilizer in butyl and bromobutyl tire innerliner formulations [49]. They reported improvements in clay dispersion and reduced air permeability. Similarly, Weng and colleagues at ExxonMobil Chemical Co. report the use of liquid maleated polymers (MAH-g-PIB, MAH-g-BIMSM) to compatibilize BIMSM–clay nanocomposites [52]. Gong et al. functionalized BIMSM polymers with maleic anhydride and then melt mixed them with organoclays [50,51]. They reported an

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RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

increase in the interlayer spacing, indicating polymer intercalation, and improvements in air impermeability ranging from 20 to 53%. 14.3.4 Low Molecular Weight Polymers and Resins Low molecular weight polymers (Mw 1,000–10,000) are widely used in the rubber industry as processing aids to reduce viscosity in rubber compounds. Examples include liquid polyisobutylene, polybutadiene, or polybutene oils and polyethylene waxes. When functionalized with polar end groups such as hydroxide, thiol, carboxylic acid, amines, epoxides, anhydrides and the like, they can serve as wetting agents to help disperse fillers in high molecular weight polymers. This concept has been applied to the dispersion of clays in rubber [61]. Polybutene process oils were investigated by Dias et al. in butyl nanocomposite tire innerliner formulations [64]. The polybutenes served to improve the processability and flexibility of the innerliner as evidenced by superior green strength and fatigue to failure properties. Air permeability properties improved, suggesting superior clay dispersion. Gong et al. reported the use of a low molecular weight amine functionalized polyisobutylene to treat Na-Mt [66]. The treated clay when used in a BIMSM tire innerliner formulation gave superior air impermeability. A similar approach is disclosed by Wang et al. using quaternary amine salts of polyisobutylene as a clay compatibilizer [69]. Elspass and Peiffer investigated a blend of low (Mw 70,000) and high (Mw 400,000) molecular weight BIMSM polymers in an OMt nanocomposite [62]. The presence of the low molecular weight polymer was reported to aid in the dispersion of the clay, as evidenced by improved mechanical and air permeability properties. In a follow-up work, Wang et al. confirmed by X-ray diffraction that adding increasing amounts of the low molecular weight BIMSM fraction helped to exfoliate the clay [68]. Hydrocarbon resins based on terpenes (polylimonene, pinene, etc.) and dicyclopentadiene have been investigated as compatibilizing agents in BIIR and BIMSM nanocomposites [55–58]. It is claimed that they improved polymer processability without detracting from the air permeability properties. This is in contrast to conventional napthenic and paraffinic processing oils, which generally result in a worsening of air permeability. Phenolic curing resins have been used in combination with butyl [60] and BIMSM [59] nanocomposites, and it is believed that these resins can also act as compatibilizing agents by promoting polymer intercalation within the clay layers. 14.4 METHODS OF PREPARATION OF BUTYL RUBBER–CLAY NANOCOMPOSITES Polymer–clay nanocomposites are principally prepared in one of four following ways: . .

melt method solution method

METHODS OF PREPARATION OF BUTYL RUBBER–CLAY NANOCOMPOSITES

. .

445

latex method in situ polymerization.

All four methods have been investigated for the production of butyl rubber nanocomposites, however, most of the literature deals with the first two (see Table 14.1). 14.4.1 Melt Method The melt method of preparing polymer–clay nanocomposites involves dispersing the clay directly within the polymer melt and can be carried out in an internal mixer, using a two-roll mill or using an extruder. This method is generally the first choice from a practical point of view since it is a straightforward one-step process that does not require any solvents or work-up. It is also most applicable for industrial applications, since most rubber processing is carried out in this way. However, in the case of nonpolar polymers such as butyl rubber, it is generally not possible to achieve good clay exfoliation by melt processing alone without the use of a polar compatibilizer agent, as discussed in Section 14.3. Bhattacharya et al. in a review of the topic, report that for effective clay exfoliation, residence time in the mixer plays a more important role than mixing shear intensity [86]. This implies that the kinetic diffusion rate for melt penetration into the clay layers is the primary factor to consider. If, for example, one is carrying out clay dispersion using a twin-screw extruder, one should incorporate into the screw design mixing elements that increase the residence time in the extruder (e.g., reversing elements). One disadvantage of incorporating clays in the melt is that once the mixing is complete and the shear stress is removed, the elastomer undergoes stress relaxation and any clay platelet alignment is lost, which is undesirable for properties such as air impermeability [37]. In contrast, when preparing nanocomposites from solution or latex, it is possible to obtain a certain degree of alignment of the clay platelets, as the solvent or water is evaporated and the clay layers lie down [87]. This is really only practical for the preparation of thin rubber coatings, as in the butyl latex technology of InMat Inc. Liang and coworkers used a modified melt intercalation process in which the clay is preswollen with an organic solvent prior to melt mixing with the rubber [78]. The method yielded butyl nanocomposites with superior properties compared to nanocomposites prepared by regular melt mixing or in solution. 14.4.2 Solution Method The solution method is carried out by adding clay to a polymer solution with vigorous agitation. Organically modified clays (OMt) that contain long alkyl groups (such as C18) will readily disperse in hydrocarbon solvents. Unmodified clays (Na-Mt), however, are too organophobic and must first be dispersed in water, then with the aid

446

RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

of an emulsifying agent and a high speed mixer, combined with the polymer solution to create an emulsion. Weng and colleagues at ExxonMobil Chemical have outlined an industrially feasible route to a BIMSM-based nanocomposite using the above approach [42–44]. The BIMSM polymer is first modified with an organic amine to form the aminefunctionalized ionomer (as shown in Figure 14.4b). Aqueous slurry of Na-Mt is then injected into the polymer cement together with a surfactant, then emulsified using a high-speed mixer. Alternately, a hydrocarbon solution of an organically modified clay may be used. The product after coagulation and drying yields a well-exfoliated clay nanocomposite with superior air barrier properties. Liang et al. conducted a comparative study of melt versus solution methods to prepare butyl–clay nanocomposites and made several observations [88]: .

.

.

The aspect ratio of the clay layers in the solution prepared butyl nanocomposite was larger than that of the melt prepared nanocomposite (Figure 14.7). This is likely due to degradation of the clay platelets under shear during melt mixing. The mechanical properties of the solution prepared nanocomposites were superior to those of the melt prepared nanocomposites. This was attributed to the fact that nanodispersed clay with higher aspect ratio has greater stress bearing capability. The larger contact surface area results in a greater interaction between the clay and the polymer, restricting the motion of the rubber chains. Gas barrier properties were superior for the solution prepared nanocomposites compared to those prepared in the melt (Figure 14.8).

FIGURE 14.7 TEM images of melt (left) versus solution prepared (right) butyl nanocomposites showing the difference in nanoclay aspect ratio. Reproduced from Ref. [88] with permission from Elsevier.

METHODS OF PREPARATION OF BUTYL RUBBER–CLAY NANOCOMPOSITES

447

1.00

PC/PP

0.95 0.90 M-IIRCN

0.85

S-IIRCN

0.80 0.75 0

2

4

6

8

Loading of clay (vol%)

FIGURE 14.8 Relative gas permeability of melt (M-IIRCN) versus solution (S-IIRCN) prepared butyl–organoclay nanocomposites. Reproduced from Ref. [88] with permission from Elsevier.

The conclusion from this study is that for the preparation of an IIR–OMt nanocomposite without the aid of a compatibilizer, the solution intercalation method gave better results than melt mixing. It should be noted however, that for the solution experiments, vigorously mixing was carried out for 24 h, while the melt blending on a two-roll mill only proceeded for 15 min. Mang and Han reported similar results in a recent study comparing melt versus solution intercalation methods [80]. The authors prepared a series of OMt clays using amine modifiers with alkyl groups increasing in chain length from C3 up to C18. As the length of the alkyl chains increased, so did the clay interlayer spacing. The mechanical properties of the IIR–OMt nanocomposites were compared. Nanocomposites prepared in solution showed consistently higher tensile strength and modulus, indicating stronger reinforcement due to better interaction between the polymer and the clay. 14.4.3 Latex Method Much of the world’s rubber is synthesized in latex form, either naturally (natural rubber, NR) or synthetically (e.g., nitrile rubber, NBR; styrene butadiene rubber, SBR; polychloroprene rubber, CR). Since natural clays (e.g., Na-Mt) are hydrophilic and readily disperse in water, it is logical that rubber–clay nanocomposites should be able to be conveniently produced in latex form. Hydration of the clay in contact with water decreases the attractive forces between the clay layers, which facilitates polymer intercalation and clay exfoliation. The latex method has been used to produce clay nanocomposites of NR [89], SBR [90–93], CR [94], NBR [95], and other rubbers.

448

RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

Butyl rubber cannot be produced by latex emulsion polymerization, since the cationic polymerization process and the initiators used are highly water sensitive. However, butyl latex can be made postpolymerization by redissolving the solid rubber in a hydrocarbon solvent and emulsifying with water-based surfactants. Butyl latex is available commercially from Lord Corporation in the United States and from PolymerLatex GmbH in Europe. A butyl latex nanocomposite technology has been developed by Goldberg, Feeney, and colleagues at InMat Inc. [29,73–77]. Vermiculite was selected as the filler of choice due to its very high aspect ratio that can be up to 10,000:1. A relatively high level of vermiculite (20–30 wt.%) can be effectively dispersed in butyl latex with the aid of surfactant wetting agents. Because the process of clay exfoliation in latex involves relatively low mixing shear, it is claimed that the large vermiculite platelets are able to survive mostly unbroken through the process, maintaining their high aspect ratio. This is in contrast to the high shear environment involved in melt blending which can degrade the size of the clay platelets [96]. Another claimed advantage of the latex process is that the clay platelets are able to align themselves as the latex film dries, creating a very efficient barrier to air diffusion. 14.4.4 In Situ Polymerization Nanocomposites may be formed in situ during a polymerization reaction if the monomers are able to intercalate within the clay layers and then undergo polymerization with the aid of an initiator. This method was first described by Blumstein for the homopolymerizations of acetonitrile and methyl methacrylate [97]. Researchers at Toyota also used this method in their early groundbreaking work on nylon nanocomposites [98,99]. The polymerization of butyl rubber proceeds via a cationic polymerization mechanism, which is very sensitive to any kind of polar impurities. It is unlikely therefore that clays could be introduced into a butyl polymerization reactor without having a detrimental effect on either the polymer conversion, molecular weight, or reactor fouling. Surprisingly, Osman and Crockett have reported that butyl polymerization reactions were successfully carried out on a laboratory scale in the presence of organoclays with only a modest 10% decrease in reaction conversion (from 80 to 70%) [100]. Furthermore, they were able to brominate the butyl nanocomposites to give the corresponding bromobutyl nanocomposites. The bromobutyl nanocomposites when compounded in a tire innerliner formulation showed a 30% improvement in air impermeability compared to the standard bromobutyl formulation, indicating that good clay dispersion had been achieved. Kim et al. reportedly polymerized isobutylene on the surface of silica nanoparticles by living cationic polymerization [101]. The resulting polymer–nanoparticle hybrids were obtained with good yield and molecular weight and were characterized by thermogravimetry, gel permeation chromatography, and dynamic light scattering measurements.

PROPERTIES AND APPLICATIONS OF BUTYL RUBBER–CLAY NANOCOMPOSITES

449

14.5 PROPERTIES AND APPLICATIONS OF BUTYL RUBBER–CLAY NANOCOMPOSITES 14.5.1 Air Barrier Properties Butyl rubber is principally used for its air barrier properties in tire innerliners and innertubes, and therefore most of the research into butyl nanocomposites has been directed toward these applications. It has long been known that addition of inorganic fillers such as clay or silica can lower the diffusion through an elastomer, since the filler generally has lower permeability to air than the polymer. However, high filler loadings must normally be used to achieve a significant effect, which can render the compounds very stiff and inelastic. Exfoliated nanoclays, however, can achieve this barrier enhancing effect at much lower filler loadings due to their very large effective surface area. This is believed to be due to the creation of a “tortuous path” that retards the passage of gas molecules through the polymer resin. Modeling studies have shown that it is theoretically possible with a fully exfoliated nanoclay to improve the barrier properties of a polymer composite by several orders of magnitude [102–104]. Table 14.2 gives a summary of permeability results for some butyl–clay nanocomposites taken from the recent literature. Results vary widely depending on (a) the type of butyl rubber and clay used, (b) whether a compatibilizer was present, and (c) the processing method used (melt, solution, or latex). All of these factors influence the degree of clay exfoliation, which is the main prerequisite to achieving excellent barrier performance. The lowest air permeability values achieved to date for a butyl nanocomposite are those associated with the butyl-vermiculite latex technology developed by Goldberg, Feeney, and colleagues at InMat Inc. In patent filings together with researchers at Michelin, the authors claim up to three order of magnitude improvement in air permeability by applying thin latex coatings to various substrates [73–75]. Takahashi et al. evaluated the technology as a barrier to various gases and reported a more modest 20- to 30-fold decrease in permeability, depending on the amount of vermiculite used in the formulation [87]. The authors showed that the experimental results could be fitted to permeability curves derived from the theoretical models, and concluded that the effective aspect ratio of the vermiculite filler was in the range of 100:1–480:1. The tire innerliner technology developed by InMat and Michelin concerns the application of a thin coating of a butyl nanocomposite latex to the inside of a vehicle tire, which upon drying results in a flexible barrier coating 20 mm in thickness, approximately 1/40th the thickness of a traditional bromobutyl innerliner. This allows for potential weight reduction of up to 1 kg/tire for a typical passenger car tire. Vulcanizing chemicals can also be added to the latex formulation so that the barrier film can be cured during the drying process. Air permeability of the nanocomposite coatings are reported to be up to three orders of magnitude lower than a typical bromobutyl tire innerliner, due to the ability of the clay particles to align themselves as the latex film dries, creating an efficient barrier to air diffusion. However, due to challenges of integrating this latex coating technology into tire

450

Surfactant wetting agents Surfactant wetting agents Vinylbenzyl chloride SBR

MAH-g-IIR BIIR phosphonium ionomer Phenolic resin

MAH-g-BIMSM

Imidazolium surfactants Liquid maleated PIB None

None

BIMSM ammonium ionomer

IIR

IIR BIIR

BIMSM

BIMSM

BIIR, BIMSM BIIR CIIR

BIMSM

BIMSM

IIR

IIR

Compatibilizer

Polymer

OMt

OMt, 3–9

Solution

Solution

39

10–40

42 0–41

Melt Melt

OMt, 12 Mica, 35–60

20–53

40–60

46–64 60

25–90

95–98

84–99.9

50

Solution, melt

Melt

Melt Melt

Melt

Latex

Latex

Processing Method

Reduction in Air Permeability (%)

Treated mica, 30

OMt, 8.5–11.5

OMt, 5–15 High aspect ratio talc, 60 OMt, 5–50

OMt, 5.5–30

Vermiculite, 17–23

Vermiculite, 10–35

Nanoclay Type and Loading (phr)

TABLE 14.2 Selected Physical Properties of Butyl Rubber–Clay Nanocomposites

–11

20–170

50 0

4–46

Increase in Tensile Strength (%)

90–200

30–100

250 0

Increase in Modulus (%)

Reduced crack growth Longer cure times (t90) Longer cure times (t90)

Longer cure times (t90)

Improved wear resistance

Other Properties

[38]

[37]

[49] [106]

[25]

[50,51]

[105]

[48] [24]

[72]

[87]

[73,75]

References

451

BIMSM ammonium ionomer None None None

Terpene hydrocarbon resins

BIIR phosphonium ionomer Phenolic resin

Phenolic resin None None

Phenolic resin

BIMSM

BIIR, BIMSM

BIIR

BIMSM

BIMSM IIR CIIR

IIR

IIR IIR IIR

None

BIIR

OMt, 5

OMt, 4 Na-Mt OMt, 0–12

OMt, 16

OMt, 18

OMt, 4

OMt, 3–7 OMt, 10 OMt, 8

OMt

HAR talc, 30

Melt

Solution Melt Solution

Solution

Melt

Melt Melt Melt; solution Melt

Solution

Melt

0

8–13

14–35 30 15; 20

36

30–38

100 20 0–10

150

550

18

120–160 100 200; 180

20–25

60 50 0–100

110

60–100 180 100; 90

Increase in tear strength Longer cure times (t90)

Shorter cure time (t90); Reduced die swell

Reduced compound viscosity

Longer cure times (t90); improved flex fatigue

[60]

[108] [88] [109]

[59]

[33]

[58]

[78] [107] [88]

[39]

[23]

452

RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

manufacturing, it has yet to see commercial use in tires. Nevertheless, researchers at InMat have gone on to successfully develop their butyl nanocomposite Air D-Fense technology for nontire applications, for example, in barrier coatings for sports balls, packaging films, and membranes [29]. A number of tire companies have issued patents related to using nanoclays to reduce air permeability in tire innerliner formulations. Fudemoto et al. (Bridgestone Corp.) evaluated treated mica clays in BIIR and BIMSM innerliner formulations and achieved up to 50% reduction in air permeability [25]. Krueger (Continental AG) added 30 phr of high aspect ratio talc to a BIIR tire innerliner formulation and achieved 30–38% reduction in air permeability [23]. Bergman (Michelin) used terpene hydrocarbon resin compatibilizers in BIMSM nanocomposite innerliner formulations and reported permeability reductions of 8–13% [58]. Maruyama et al. (Yokohama Rubber Co.) used a low molecular weight liquid maleated polyisobutylene (MAH-g-PIB) as a clay compatibilizer in butyl and bromobutyl tire innerliner formulations and reported up to 42% reduced air permeability [49]. Song (Kumho Tires Co.) used a vinylbenzyl chloride functionalized SBR to compatibilize montmorillonite clays with butyl rubber and reported permeability reductions as high as 90% [72]. Miyazaki (Sumitomo Rubber Industries Ltd.) used high loading of mica nanofiller in CIIR/NR/BR innerliner formulations to reduce the air permeability by up to 41% [106]. Tsou and Measmer prepared BIIR and BIMSM nanocomposites with varying amounts of OMt clay and observed up to 40% reduction in air permeability [37]. The authors reported a good correlation between the experimental results and predicted permeability from Gusev–Lusti equations. From TEM and X-ray imaging and analysis, BIMSM polymers reportedly gave better clay dispersion than BIIR and exhibited correspondingly lower air permeability. Zhang et al. studied the effect of curing temperature on the permeability properties of sulfur-cured IIR–OMt nanocomposites [107]. As the curing temperature increases above 150 C, the permeability properties worsen (Figure 14.9). This was said to be due to the quaternary amine clay modifiers starting to react with the sulfur curatives, resulting in desorption from the clay surface and reagglomeration of the clay layers. Kato and colleagues used low molecular weight MAH-g-IIR as a compatibilizer for butyl rubber and OMt clay and reported up to two-thirds reduction in air permeability for the resulting composites [48]. Liang et al. used a modified melt intercalation process in which the clay is preswollen with an organic solvent prior to melt mixing with butyl rubber [78]. Up to 35% reduction in air permeability was observed compared to the base polymer. 14.5.2 Reinforcement Properties The reinforcing effect of exfoliated clays has been documented for butyl rubber in numerous studies as shown in Table 14.2. Some representative examples are described below. Parent et al. [33] observed a dramatic increase in polymer reinforcement when a bromobutyl phosphonium ionomer was melt compounded with OMt clay. As shown

PROPERTIES AND APPLICATIONS OF BUTYL RUBBER–CLAY NANOCOMPOSITES

453

1.0

PC/PP

0.9 0.8 0.7

0.6 0.5 110

120

130

140

150

160

170

180

Curing temperature (ºC)

FIGURE 14.9 IIR–OMt nanocomposite gas permeability as a function of curing temperature. Reproduced from Ref. [107] with permission from John-Wiley & Sons.

in Figure 14.10, as the clay content is increased from 0 to 15 wt.%, the modulus of the uncured ionomer nanocomposite increases by over fourfold and the slope of the stress–strain curve became sharply steeper. The reinforcing effect for bromobutyl rubber in the absence of ionomer was not as significant. Maiti et al. prepared a series of BIMSM–montmorillonite nanocomposites in solution and evaluated their mechanical properties [59]. Using organically modified montmorillonite clays, a substantial increase in reinforcement was seen as the clay loading increased. With 16 phr of OMt, the tensile strength of the composite was 2.5 times that of the base polymer. When Na-Mt was used, little reinforcement was seen, indicating relatively poor compatibility with the polymer. In a related study, the authors studied the effect of OMt clays in carbon black filled BIMSM compounds and observed a synergistic reinforcing effect of the clay with the carbon black filler [108]. Sridhar et al. recently documented a similar synergistic reinforcing effect between carbon black and OMt clay in chlorobutyl compounds [109]. Butyl rubber is used in tire curing bladders that require the polymer to withstand severe service conditions of heat and pressure, and multiple cycles of expansion and contraction. As a result, curing bladder life tends to be short and they must be replaced often, which results in unwanted process down time. Bladder failures occur due to tears and rupture, and therefore a need exists to improve the strength of the compounds to resist such failures. Samadi and Kashani completed a study comparing various OMt clays in a resin cured curing bladder formulation [60]. The clays were commercially sourced and differed only in the quaternary ammonium ions used. Cloisite 10A which is modified with hydrogenated tallow benzyl dimethyl ammo nium ions, showed the largest interlayer spacing of 43 A when melt compounded with butyl rubber, suggesting that the polymer chains were well intercalated within the

454

RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

9.0 IIR–PPh3Br/NR4+ –MM

8.0

5 wt. % E=0.79 MPa

Stress (MPa)

7.0 15 wt. % E=2.15 MPa

6.0

3 wt. % E=0.64 MPa

5.0 4.0

Unfilled E=0.55 MPa

3.0 2.0 1.0

Uncured BIIR

0.0 3.5 BIIR–ZnO/NR4+ –MM

Unfilled E=0.33 MPa

Stress (MPa)

3.0 2.5

5 wt. % E=0.43 MPa

2.0 1.5

15 wt. % E=0.65 MPa

1.0 0.5 0.0 0

20

10

30

40

Strain (mm/mm)

FIGURE 14.10 Stress–strain curves of IIR–PPh3Br and BIIR nanocomposites of NR4 þ –Mt. Reproduced from Ref. [33] with permission from Elsevier.

clay layers. It was postulated that the phenolic curing resin in the formulation assisted in the intercalation process. Elastic modulus was seen to increase by 20–25% over the base compound. 14.5.3 Vulcanization Properties Mineral fillers are known to alter the vulcanization properties of halogenated butyl rubber compounds depending on their pH as well as their moisture content [1]. Acidic fillers such as calcined clay, have an accelerating effect, while alkaline fillers and those that are hygroscopic in nature, tend to retard the cure. Nanoclays that have been

PROPERTIES AND APPLICATIONS OF BUTYL RUBBER–CLAY NANOCOMPOSITES

455

organically modified with quaternary amines become slightly alkaline and therefore fall into the second category, exhibiting a retarding effect on halobutyl rubber vulcanization. The onset of cure (time to scorch) is also delayed, which may be beneficial in applications where additional scorch protection is required. Note that the opposite effect is seen for other polymers: cure acceleration and reduced scorch times have been observed with organoclays in NR, SBR, and EPDM compounds [110]. Rodgers et al. observed an increase in 90% cure time (t90) when an amine modified BIMSM nanocomposites was used in a sulfur cured innerliner compound [38]. Cure time, t90 increased by 50% for the nanocomposite compound compared to the standard innerliner formulation. It was found however that the retarding effect could be compensated for by adjusting the stearic acid content in the formulation. Samadi and Kashani similarly observed increased cure and scorch times in resin cured butyl bladder formulations when organoclays were introduced [60]. Soisson et al. investigated the cure properties of a BIMSM nanocomposite in a tire innerliner formulation [105]. The cure rate was significantly slower compared to a standard bromobutyl compound, and therefore different cure accelerators were evaluated. Dibutyl thiourea (DBTU) was found to be an effective accelerator for overcoming the cure retarding effect of the clay. A shorter scorch time and faster cure was found to be necessary to prevent reagglomeration of clay particles. Quite a bit of work has been done to look at what happens to clay morphology during the curing process. During the prescorch period before vulcanization kicks off, the polymer chains and clay platelets are mobile and can rearrange, thus altering the nanocomposite morphology. Usuki and Kato observed increased clay exfoliation after curing an EPDM nanocomposite and attributed this to the clay layers being pulled apart as the polymer chains undergo cross-linking reactions [111]. Interestingly, Varghese and Karger-Kocsis observed the opposite effect with natural rubber nanocomposites, namely a collapse of the clay interlayer spacing after vulcanization [112]. Liang and coworkers studied the effects of curing temperature and pressure on butyl rubber nanocomposite morphology [113–115]. IIR–OMt nanocomposites were prepared on a two-roll mill at 30–60 C in the absence of curatives. The authors then compared the clay morphology, first on the untreated samples at room temperature, then after heating for 1 h at 160 C under atmospheric pressure in an oven, and finally after heat treatment under high pressure in a curing press. After atmospheric pressure heat treatment, the clay interlayer spacing shrunk and clay aggregation was seen in the TEM imaging. When the heating time was shortened to 10 min, the effect was not as pronounced. These findings are consistent with what is generally known in the literature concerning filler reagglomeration. Heating under high pressure resulted in polymer molecules being extruded out from between the clay interlayer spaces, speeding up the process of clay reagglomeration. From these observations, the authors made the following recommendations for curing of polymer nanocomposites. Select a curing chemistry that kicks-off at lower temperatures, with a short scorch time and a fast rate of cure. These curing conditions will help to minimize the extent of clay reagglomeration.

456

RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

Zhang et al. investigated the effect of curing temperature on the microstructure and properties of IIR–OMt nanocomposites [107]. They found that curing temperature strongly influences the final state of clay dispersion as seen by TEM and X-ray analysis. If the curing temperature is too low, the longer cure time results in the clay layers reagglomerating. If the curing temperature is too high, the quaternary amine clay modifier can desorb from between the clay interlayer spaces similarly causing the clay layers to contract. A curing temperature of 160 C was found to be close to ideal. Curing temperature was also found to affect the permeability and tensile properties of the nanocomposites (see Figure 14.9). 14.5.4 Adhesion Properties An important application area for butyl rubber is in adhesives and sealants, and therefore improving the adhesion properties of the polymer could be of significant interest. Osman prepared a maleic anhydride-grafted bromobutyl (MAH-g-BIIR) nanocomposite and investigated its adhesive properties [54]. The MAH grafting reaction was carried out in solution using a novel metal alkanoate catalyst, which resulted in very little molecular weight breakdown. MAH grafting could be carried out in the presence of OMt clay to give a MAH-g-BIIR nanocomposite in a one-step process. The adhesive strength of the maleated bromobutyl and its nanocomposite were measured in conjunction with various substrates (Table 14.3). As expected, maleation of the polymer increased the bonding ability to polar substrates. Surprisingly, the nanocomposite showed even greater adhesion, which is a significant finding since adhesion often suffers when compounding with clay fillers. It can be postulated that the reinforcing effect of the nanoclay serves to increase the cohesive strength of the elastomer. Adhesive tack is the ability of two unvulcanized elastomer surfaces to adhere together when in contact for a short period of time under light pressure. This characteristic is important in industrial rubber applications such as tire building where the butyl innerliner must have sufficient tack to adhere to the adjacent carcass ply layers. Bhowmick et al. investigated the effect of sepiolite clay on the autohesive tack of BIMSM rubber [30]. With 2 and 8 phr of sepiolite clay in the formulation, the adhesive tack strength of the rubber increased by 140 and 300%, respectively compared to the base polymer. The elastic modulus was also enhanced. The authors posit that due to the fine fibrous nature of the clay particles, the polymer molecules are TABLE 14.3 Adhesive Strength of an MAH-g-BIIR–Clay Nanocomposite [54] Adhesive Strength (psi) (Tel-Tak Test)

Stainless steel Glass Mylar Teflon

BIIR

MAH-g-BIIR

MAH-g-BIIR þ 7 phr Nanoclay

11.3 6 2 1.5

33.3 30.8 9.7 6.5

51 35.3 30 10.7

CONCLUSIONS

457

able to reach across the particles and maintain a well-entangled network, while at the same time benefiting from the reinforcing effects of the filler. 14.5.5 Other Properties Processability is an important factor in most butyl rubber applications, such as in the critical extrusion and calendering steps in the tire-building process. Poor processability can adversely impact the economics of a manufacturing process due to reduced production rates and an increase in off-specification product. Lohse reports that BIMSM nanocomposites in innerliner formulations allow for an increase in the extrusion rate during processing compared to the standard formulation [116]. Maiti et al. observed a decrease in die swell when BIMSM–clay nanocomposites were compared to the parent polymer [59]. The improvement was attributed to better polymer–filler interactions, reduction in elastic nature, and development of higher shear viscosities. Dias et al. investigated the use of polybutene oils to further improve the processability of BIMSM nanocomposites in tire innerliner formulations [64]. Sridhar et al. studied the effect of OMt clay on the relaxation behavior of chlorobutyl rubber, comparing conventionally filled compounds using carbon black and silica versus clay filled systems [109]. The nanoclays were found to have a strong influence on the relaxation behavior due to enhanced polymer–filler interaction. Rodgers et al. looked at the effect of aging on the barrier performance of a typical bromobutyl tire innerliner compound vs. a BIMSM nanocomposite innerliner [117]. They observed that the nanocomposite innerliner was more stable to accelerated hot air oven aging than the reference innerliner compounds and the barrier properties were more consistently maintained over time. The authors predicted better long-term tire performance using a nanocomposite innerliner.

14.6 CONCLUSIONS Nanoclays continue to be the subject of intense research interest in the rubber industry for their potential to significantly enhance elastomer performance, particularly with regard to mechanical and air barrier properties. Yet despite all the recent activity, an industrial rubber application has yet to be commercialized that could realize the potential of nanoclay technology on a large scale. Butyl rubber, because it is so heavily used in barrier applications, could turn out to be the ideal elastomer to exploit the commercial potential of nanoclays. The vast majority of research into butyl–clay nanocomposites is currently being carried out in corporate research labs directed toward the tire industry, which gives an indication of the focused commercial interest in this technology. Current research efforts into butyl–clay nanocomposites have concentrated on ways to improve the compatibility of the elastomer with the clay, since it is inherently difficult to achieve good clay dispersion in nonpolar polymers such as butyl rubber. The main compatibilization techniques explored to date include:

458 . .

.

RUBBER–CLAY NANOCOMPOSITES BASED ON BUTYL AND HALOBUTYL RUBBERS

The use of surfactants or swelling agents to further separate the clay layers and aid polymer intercalation. Modification of the butyl polymer chains by the addition of polar functional groups to increase the affinity with the clay surface. Examples include butyl rubber ionomers, and maleic anhydride grafted butyl rubber. The use of suitable low molecular weight polymers and resins, which can intercalate within the clay layers while being compatible with the butyl rubber matrix.

Most of the known methods of dispersing clays into polymers have been explored with butyl rubber, including melt-mixing techniques, solution and latex methods, as well as in situ incorporation during polymerization. Melt-mixing methods are the most industrially practical, but in general, solution and latex methods have been found to be more effective for clay exfoliation, resulting in superior physical properties. Although the barrier enhancing effect will continue to be the main reason for the interest in using nanoclays with butyl rubber, research into the mechanical properties of the composites suggests that there will also be benefits in applications that require resistance to mechanical and environmental stresses. More research needs to be done into the effects of nanoclays on the rheological properties of butyl rubber, for example, the effects on damping properties and hysteresis, which are important characteristics of the polymer. Questions related to vulcanization chemistry, aging properties, and processability need to be further explored as the science of butyl–clay nanocomposites evolves from the research lab toward practical industrial uses.

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CHAPTER 15

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM) MAKOTO KATO

15.1 INTRODUCTION Ethylene propylene rubber (EPR) is a widely used synthetic polyolefin rubber that has great commercial importance. EPR is a copolymer in which the E refers to ethylene, P to propylene, and R to rubber. Two different kinds of synthetic rubber are used to produce EPR—EPM and EPDM. The M in these abbreviations refers to a classification in ASTM standard D-1418. “M” class rubbers possess a saturated chain of the polymethylene type. The D stands for diene that is introduced into the rubber so that it can be vulcanized. EPM (ethylene propylene monomer (M-class) rubber) is a copolymer which has excellent insulation, water resistance, and UV-resistance properties. It is used as insulation for high voltage cables and as a raw material for automotive interior applications. In EPDM rubber (ethylene propylene diene monomer (M-class) rubber, a small amount of a third monomer (the diene) is introduced to provide improved heat resistance compared to an EPM rubber. Typical dienes include DCPD (dicyclopentadiene), ENB (ethylidene norbornene), and VNB (vinyl norbornene). ENB is currently most widely used in the production of EPDM because it affords the fastest vulcanization rate. EPDM is commonly vulcanized with sulfur or peroxide. EPDM is easily vulcanized in comparison with EPM, has excellent electrical properties, and exhibits superior heat, ozone, and weather resistance. EPDM is the most widely used synthetic rubber in a wide variety of applications ranging from automotive hoses, belts and weatherseals to gaskets, roof sheeting and glazing seals to wire and cable insulation. Use of clay minerals in EPR applications provides benefits such as improved insulation properties, enhanced gas and water impermeability, and reinforcement for Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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greater mechanical properties. Carbon black is most widely used for reinforcement of EPR. A large amount of this material is necessary, though, to achieve the desired level of reinforcement and other objective properties. If the same effects can be achieved using a small amount of clay minerals, reduction in both the size and weight of automotive parts is expected. This chapter separately introduces EPDM–clay nanocomposites and EPM–clay nanocomposites from the viewpoint that the preparation methods and the characteristic are greatly different whether to vulcanize.

15.2 TYPES OF CLAY MINERALS USEFUL IN EPM–, EPDM–CLAY NANOCOMPOSITES Three types of clay minerals are used in EPM– and EPDM–clay nanocomposites: montmorillonite, fluorohectorite, and Mg–Al layer double hydroxide (LDH). The chemical formula of montmorillonite is (Na,Ca)0.33(Al,Mg)2(Si4O10) (OH)2
nH2O. Na þ and Ca þ cations that are exchangeable with organic onium ions such as ammonium, phosphonium, and sulfonium ions, giving the montmorillonite cation exchange capabilities. Ammonium ions (primary, secondary, tertiary, or quaternary) are easy to obtain and thus are usually used as organic modifiers. In the preparation of EPM– or EPDM–clay nanocomposites, many researchers including Usuki and coworkers and many other researchers use hydrophobic organic modifiers such as octadecyl ammonium (C18H37NH3 þ ) [1–5], octadecyl trimethyl ammonium {(C18H37)(CH3)3N þ } [6,7], dioctadecyl dimethyl ammonium {(C18H37)2(CH3)2N þ } [8–10], and hexadecyl ammonium (C16H33NH3 þ ) [11–14]. Zheng and coworkers tried to prepare a semitransparent EPDM–clay nanocomposite by using methyl bis(2-hydroxyethyl)octadecyl ammonium {CH3C18H37(CH2CH2OH)2N þ } as an organic modifier [15]. Zheng also investigated octadecyl trimethyl ammonium, methyl bis(2-hydroxyethyl)octadecyl ammonium, and dimethyl benzyl octadecyl ammonium {(CH3)2C6H5CH2C18H37N þ } [16] in order to examine the influence of the chemical structure of organic modifiers on the properties of rubber nanocomposites. Gatos and Karger-Kocsis used octadecyl ammonium and octadecyl trimethyl ammonium [17]. Alternatively, Jeon and coworkers tried to disperse Na-montmorillonite (without ammonium) into rubbers including EPM by using polyethylene glycol (PEG) as a modifier [18]. Wu and coworkers used fluorohectorite, a synthetic 2:1 silicate with the chemical formula Na0.33Mg2.67Li0.33Si4O10F2
nH2O, to examine the effect of the different characteristics of EPDM, NR, and SBR rubbers on the structure and properties of rubber–clay nanocomposites [19]. In this study, octadecyl ammonium was employed as the organic modifier. Acharya and coworkers synthesized EPDM–LDH nanocomposites [20]. LDHs (hydrotalcite-like systems) are layered materials consisting of positively charged layers, which are in contrast to cationic clay minerals that possess negatively charged layers. LDHs can be represented by an ideal chemical formula of [MII(1–x)MIIIx(OH)2]x þ Am  x/m/nH2O], where MII is a divalent metal ion

COMPATIBILIZER SYSTEMS FOR OLEFINIC RUBBER–CLAY NANOCOMPOSITES

467

(i.e., Mg2 þ ), MIII is a trivalent metal ion (such as Al3 þ ), and A is an exchangeable interlayer anion. Because an LDH is a 1:1 type layered material, the layers are thinner than those of montmorillonite and are typically 0.48 nm thick. Acharya et al. used a Mg–Al LDH with CO3 between the layers to prepare EPDM/Mg–Al LDH nanocomposites by exchanging the CO3 with dodecyl sulfate (C12H25SO4 ).

15.3 COMPATIBILIZER SYSTEMS FOR OLEFINIC RUBBER–CLAY NANOCOMPOSITES EPM– and EPDM–clay nanocomposites cannot be prepared easily. The polyolefin rubbers are hydrophobic materials and thus have poor miscibility with hydrophilic clay minerals such as montmorillonite. When preparing a polypropylene (PP)–clay nanocomposite, Usuki et al. reported that the dispersibility of clay minerals in a PP matrix was improved, but PP–clay nanocomposites were not obtained even if montmorillonite was used with the highly hydrophobic organic modifier dioctadecyl dimethyl ammonium. A PP–clay nanocomposite was obtained for the first time by using polyolefin oligomers containing hydroxyl, carboxyl, maleic anhydride, and other polar groups as compatibilizers [21–23]. Since that initial report, compatibilizers have been commonly employed in systems that exhibit poor miscibility. Gatos et al. studied the effects of using grafted glycidyl methacrylate as a compatibilizer for EPDM–organoclay nanocomposites [4]. They prepared an EPDM rubber containing grafted glycidyl methacrylate (EPDM–GMA) groups having a GMA content of 8.7%. Clay nanocomposites made from EPDM/ EPDM–GMA (50/50 wt.%) and EPDM alone, each with 10 phr of octadecyl ammonium treated montmorillonite was prepared. These nanocomposites were vulcanized using sulfur with the different accelerators N-cyclohexyl-2-benzothiazole sulfenamide (CBS), 2-mercaptobenzothiazole (MBT), and zinc diethyldithiocarbamate (ZDEC). Addition of the grafted GMA increased the tensile strength of the rubber nanocomposites. Specifically when ZDEC was used as an accelerator, noteworthy effects were observed. The modulus at 100, 200, and 300% strain nearly doubled while still retaining a high value for strain at break compared to the values obtained for the nanocomposite without grafted GMA. These results were attributed to the formation of Zn complexes of the curatives and the octadecyl ammonium intercalant, which contributed to both enhanced intercalation/exfoliation phenomena and reaggregation of the layers [3,24]. This interaction could be seen in the XRD spectra. The basal spacing of the initial organoclay (2.10 nm) was increased in the EPDM–clay nanocomposite to 3.28 nm. When the compatibilizer EPDM-g-GMA was used, two distinct peaks were observed. The first peak appeared at 2.2 (3.92 nm) and the second peak was a broad peak in the region of 2u ¼ 5–7 (1.77–1.29). The first peak was associated with the intercalated montmorillonite with EPDM or EPDM-g-GMA, while the second peak, which was lower than the initial basal spacing of the organoclay, suggested some reaggregation of the silicate layers. Gatos et al. speculated that this confinement was due to the formation of a Zn complex in which sulfur and the amine group of octadecyl ammonium participated. Accordingly, when

468

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

octadecyl ammonium was extracted from the interlayer regions, the layer collapsed and reaggregated. Mohammadpour and Katbab studied EPDM–clay nanocomposites in which maleic anhydride-grafted to the EPDM (EPDM-g-MAH) served as an interfacial compatibilizer [9]. They reported that the content of EPDM-g-MAH influenced the dispersion status of the clay mineral (montmorillonite treated with dimethyl dehydrogenated tallow (mainly C18H37) ammonium). As the compatibilizer content increased, the dispersion status improved. The researchers concluded that an almost complete dispersion was obtained when the compatibilizer–clay ratio was 3. They also investigated the effects of EPDM molecular weight and its structural branching on the type of nanocomposite microstructure formed during a shear melt-mixing process. Three EPDM rubber grades with different structural characteristics were used: Keltan 2340A (DSM), Keltan 8340A (DSM), and Vistalon 7500 (Exxon Mobil). The ethylene content of these EPDMs was around 55% and all three incorporated ethylidene–norbornene as the diene. The Mooney viscosity (melt viscosity) and linearity of the polymer chains, however, were different. Keltan 2340A is a low Mooney viscosity (25: ML1 þ 4 at 125 C, low molecular weight) EPDM with controlled long-chain branching, Keltan 8340 is a high Mooney viscosity (85, high molecular weight) EPDM with controlled long-chain branching. Vistalon 7500 is a high Mooney viscosity (85, high molecular weight) linear chain EPDM. The shift of the sharp d001-plane spacing peak of the organoclay from 2u ¼ 5 (3.4 nm) to the lower diffraction angles in the three EPDM–EPDM-g-MAH–clay nanocomposites indicated that EPDM was able to intercalate into the interlayer spaces of the organoclay in the presence of EPDM-g-MAH as a compatibilizer, whereas no significant change occurred in the absence of the compatibilizer. The d001 spacing of the nanocomposite based on Keltan 2340A was shifted to a lower angle along with increasing lower peak density. This result provided evidence of greater interdiffusion of the low molecular weight EPDM chains and therefore more intercalation of the silicate layers. In comparing the XRD pattern of the nanocomposite based on the high Mooney viscosity and linear EPDM (Vistalon 7500) and that of the nanocomposite based on the high Mooney viscosity and branched EPDM (Keltan 8340), it can be seen that the d001 spacing of the former exhibited a lower angle with a decrease in the peak intensity compared with that of the latter nanocomposite. This difference was attributed to the steric hindrance effect of the branches, which retarded the diffusion of the chains into the interlayer spaces of the organoclay, leading to the lower extent of interaction and therefore less delamination of the clay silicate layers. Gatos and Karger-Kocsis studied the effects of primary amine (PRIM) and quaternary amine (QUAT) intercalants on the organo montmorillonite in an EPDM/EPDM-g-MA (50/50) system with sulfur and ZDEC as the curing agent [17]. They concluded that the dispersion stage of silicate layers of PRIMMt was better than that of QUAT-Mt from XRD and transmission electron microscopy (TEM) data. When PRIM-Mt was used as an organoclay, an exfoliated nanocomposite was obtained. On the other hand, when QUAT-Mt was used, an intercalated nanocomposite was obtained. They speculated that the primary amine of PRIM-Mt created a complex with the vulcanization curatives during curing.

PREPARATION OF EPDM–CLAY NANOCOMPOSITES

469

Nieuwenhuizen et al. also studied the actions of amines with zinc accelerator complexes in sulfur-vulcanized systems [25,26]. They suggested a mechanism whereby nucleophilic attack of an amine on the carbon atom of the thiocarboxy group of a bis(dialkyldithiocarbamate)zinc(II) (ZDAC) complex yields an aminedithiocarbamic intermediate from which, in the case of primary amines, thiourea products are obtained [27]. Karger-Kocsis et al. proposed that such a complex can accelerate the delamination of the clay silicate layers. Acharya and coworkers studied EPDM–clay nanocomposites in which ethylene vinyl acetate copolymer (EVA-45) with 45% vinyl acetate content was used as an interfacial compatibilizer [13]. These nanocomposites contained 2–8 wt.% of montmorillonite treated with hexadecyl amine (16Me-Mt) and were prepared by a solution intercalation method in which EPDM, EVA-45, and 16Me-Mt were mixed in dry toluene followed by curing with dicumyl peroxide under compression at 145 C for 45 min. They were found to possess a number of stacked montmorillonite and individual clay layers. Acharya considered these nanocomposites to be a mixed intercalated–exfoliated nanocomposite.

15.4 PREPARATION OF EPDM–CLAY NANOCOMPOSITES BY AN IN SITU INTERCALATION METHOD Kawasumi and coworkers pointed out that the extent of dispersion of the silicate layers of a clay in a polypropylene–clay nanocomposite depends on the miscibility of the matrix and the compatibilizer. In order to enhance the fine dispersion state of silicate layers of a clay, it is necessary to use a modified polymer matrix with the same composition as the compatibilizer. However, it is very difficult to modify EPDM rubbers because the cross-linking reaction occurs during modification process. Miscibility of the compatibilizer in a polymer matrix is also dependent on the amount of polar groups in the compatibilizer. To address this issue, Usuki et al. prepared an EPDM–clay nanocomposite using an EPDM vulcanized with modifying accelerators [1]. Five types of vulcanization accelerators for in situ modification were investigated: thiourea (ethylenethiourea, NPV/C), thiazole (2-mercaptobenzothiazole, M), sulfenamide (N-cyclohexyl-2-benzothiazylsulfenamide, CZ), thiuram (tetramethylthiuram monosulfide, TS), and dithiocarbamate (zinc dimethyldithiocarbamate, PZ). EPDM and montmorillonite treated with octadecyl ammonium (C18H37NH3 þ ) (C18-Mt) were mixed at 200 C in a twin-screw extruder. The obtained blended material zinc oxide (5 phr), stearic acid (1 phr), sulfur (1.5 phr), and the vulcanization accelerator (1.5 phr) were compounded by using a mixing roll. The vulcanized EPDM was then press molded at 160 C for 30 min to yield EPDM–clay nanocomposites. The XRD patterns of C18-Mt and EPDM–clay nanocomposites using the five types of vulcanization accelerators are shown in Figure 15.1. In the case of the EPDM–clay nanocomposites prepared using NPV/C, M, and CZ, there is a peak at a diffraction angle between 2 and 5 that clearly indicates the formation of an intercalation type of EPDM–clay nanocomposite in which the silicate

470

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

FIGURE 15.1 XRD patterns of (a) C18-Mt and (b–f) EPDM–clay nanocomposites using five types of vulcanization accelerators: (b) ethylenethiourea (NPV/C); (c) 2-mercaptobenzothiazole (M); (d) N-cyclohexyl-2-benzothiazylsulfenamide (CZ); (e) tetramethylthiuram monosulfide (TS); (f) zinc dimethyldithiocarbamate (PZ).

layers of montmorillonite are stacked, but the interlayer distance is larger than that of C18-Mt. In the case of TS and PZ, however, the peak generated by the orderly stacked layers of C18-Mt disappeared and the diffraction strength of the nanocomposite gradually shifted toward a lower angle. This diffraction pattern indicates the formation of the exfoliation type of EPDM–clay nanocomposite. From the TEM observation, the clay silicate layers were exfoliated and dispersed uniformly in the EPDM as monolayers when using the PZ vulcanization accelerator. A proposed mechanism of EPDM intercalation into an interlayer space of montmorillonite was developed by evaluating the XRD patterns taken during various stages of the EPDM–clay nanocomposite process using PZ as the vulcanization accelerator (Figure 15.2). The XRD pattern obtained after the EPDM and C18-Mt were melt compounded using a twin extruder is shown in (b). The XRD pattern in (c) was obtained after the first mixture and the vulcanization agent were mixed in a mixing roll. The final XRD pattern shown in (d) was obtained after vulcanization at 160 C. After vulcanization, the peaks at lower angles representing orderly stacked layers of clay disappeared. From these results, the researchers concluded that the intercalation of EPDM molecules into the interlayer spaces of montmorillonite occurred during the mixing roll step, where the modified EPDM was formed and intercalated into the interlayer spaces of montmorillonite, and the vulcanization process, where the clay was

PREPARATION OF EPDM–CLAY NANOCOMPOSITES

471

FIGURE 15.2 XRD patterns of (a) C18-Mt and (b–d) each stage of the EPDM–clay nanocomposite preparation process: (b) after EPDM and C18-Mt are melt compounded using an extruder; (c) after EPDM and C18-Mt, sulfur, ZnO and vulcanization accelerator PZ are compounded using a mixing roll; (d) after vulcanization at 160 C.

exfoliated and the EPDM cured. Thiuram (TS) and dithiocarbamate (PZ) type vulcanization accelerators dissociate into radicals themselves when the temperature is raised. The radicals combine with carbon atoms in the EPDM chains to polarize the EPDM molecules. These EPDM molecules intercalate into the interlayer spaces of montmorillonite through hydrogen bonds between the polar substituents on the EPDM and the clay surface. The conceptual diagram of an EPDM–clay nanocomposite is illustrated in Figure 15.3. Usuki et al. noted that the properties of the exfoliation-type nanocomposites that were obtained with the TS or PZ vulcanization accelerators were superior to those of the intercalation-type nanocomposites that were obtained with the NPV/C, M, and CZ accelerators. Tensile strength, elongation at break, tensile stress at 100% elongation, storage modulus, and gas barrier properties of the exfoliation-type nanocomposites were higher than those of the intercalation type. Zheng and coworkers studied the influence of clay modification on the structure of EPDM–clay nanocomposites prepared using the in situ method [16]. Three different montmorillonites were employed in the study: trimethyloctadecyl amine modified montmorillonite (C18a-Mt), methyl bis(2-hydroxyethyl)coco alkylamine modified montmorillonite (C12-Mt), and dimethyl benzyl octadecylamine modified montmorillonite (C18b-Mt). They evaluated the dispersibility of these three different organoclays in EPDM–organoclay nanocomposites vulcanized with a dithiocarbamate (PZ) type vulcanization accelerator system. After obtaining a melt compound,

472

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

FIGURE 15.3 Conceptualization of the vulcanization process for EPDM–clay nanocomposites and intercalation of EPDM into the interlayer space of the clay.

the EPDM samples were cured and the dispersion state of the silicate layers investigated by XRD and TEM. When C12-Mt was used as the organoclay, an exfoliated nanocomposite was obtained. On the other hand, when C18a-Mt or C18bMt was used, intercalated nanocomposites were obtained. Zheng and coworkers attributed the differing results to the presence of polar groups in C12-Mt that have a different aliphatic nature compared to EPDM. Consequently, the system was not at theta conditions and there was a favorable excess enthalpy to promote C12-Mt dispersion in the EPDM matrix. They concluded that the reinforcing ability of these

CHARACTERISTICS OF EPDM–CLAY NANOCOMPOSITES

473

various fillers in an EPDM matrix follows the order: Na-Mt (not modified montmorillonite) < C18a-Mt < C18b-Mt < C12-Mt.

15.5 CHARACTERISTICS OF EPDM–CLAY NANOCOMPOSITES 15.5.1 Gas Barrier Properties of EPDM–Clay Nanocomposites An important feature of polymer–clay nanocomposites is their excellent gas barrier properties. This gas barrier effect can be explained by the formation of a torturous diffusion path resulting from the presence of the dispersed silicate layers of clay. The permeability coefficient of a gas traveling through a polymer–clay nanocomposite can be analyzed using a geometrical model in which the silicate layers are dispersed. In films of polymer–clay nanocomposites, silicate layers are aligned nearly parallel with the film surface. According to Nielsen, the diffusion coefficient D of a liquid or a gas can be calculated if the liquid or gas is in a composite material in which plate-like particles are in a planar orientation using the following equation [28]: D ¼ D0 =f1 þ ðL=2dÞVg where D0 is the diffusion coefficient in a matrix, L is the size of one side of a plate-like particle, d is the thickness, and V is the volume fraction of particles. When montmorillonite is used at a 2% level, L is 100 nm, d is 1 nm, and V is 0.0074, and thus D/D0 is calculated to be 0.73. In the nylon 6–clay hybrid (NCH), the calculated gas barrier effect D/D0 is actually equivalent to the experimental value obtained for hydrogen (0.70) and water (0.63) [29]. Chang evaluated oxygen permeabilities of two types of EPDM–clay nanocomposites [2]. One was prepared from EPDM and octadecyl amine treated montmorillonite (C18-Mt). Another type was prepared by using low molecular weight (liquid) EPDM as a compatibilizer between EPDM and C18-Mt. These nanocomposites were vulcanized by sulfur with CBS (N-cyclohexyl benzotriazyl sulfonamide) as a vulcanization accelerator. In EPDM–clay nanocomposites filled with 10 phr C18-Mt, the oxygen permeability of the EPDM/C18-Mt system was decreased to 78% and that of EPDM/liquid EPDM/C18-Mt system was decreased to 60% relative to unfilled EPDM. Usuki and coworkers reported nitrogen permeabilities of EPDM–clay nanocomposites prepared via the in situ intercalation method (Table 15.1) [1]. For EPDM filled with 7 phr of C18-Mt and vulcanized with sulfur and PZ, the nitrogen permeabilities decreased to 70% relative to the values obtained for EPDM with no clay. Mohammadpour and Katbab studied the oxygen permeability of EPDM–clay nanocomposites prepared from low and high molecular weight EPDM rubbers with different chain linearities [9]. EPDM–clay nanocomposites prepared using high Moony viscosity EPDM showed lower permeabilities. In particular, EPDM–clay nanocomposites based on high molecular weight EPDM and 15 wt.% EPDM-g-MAH as a compatibilizer were found to have permeability levels decreased to 54% of the

474

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

TABLE 15.1 Properties of EPDM–Clay Nanocomposites with Different Vulcanization Accelerators System

EPDM–ClayNanocomposite(C18-Mt7 phr)

Vulcanization Accelerator

NPV/C

Properties Tensile strength (MPa) Elongation at break (%) 100% tensile stress (MPa) Storage modulus (MPa, 10 Hz, 25 C) N2 permeability at 60 C (10  9 cm3 cm/(cm2 s cmHg))

EPDM

M

PZ

CZ

TS

PZ

8.0 307 1.9 5.5

7.5 173 1.8 5.3

8.1 185 1.7 5.8

11.0 443 2.3 6.0

10.1 520 2.3 6.2

5.0 280 1.6 3.3

1.8

2.2

2.1

1.7

1.7

2.4

level observed for high molecular weight EPDM polymers without added clay. They concluded that the silicate layers of the organoclay were dispersed to a greater extent as a result of high shearing imposed on the silicate layers during melt mixing in a Banbury-type internal mixer. 15.5.2 Rheological Properties of EPDM–Clay Nanocomposites The rheological properties of EPDM–clay nanocomposites were investigated to determine their microstructure or the dispersion state of the silicate layers. Mohammadpour and Katbab analyzed the dynamic melt rheology behavior for different prepared EPDM–clay nanocomposites prepared from low and high molecular weight EPDM polymers with different chain linearities [9]. Values for the storage and loss modulus and relaxation time distribution function were compared with those of corresponding neat EPDM polymers. The dispersion state of the silicate layers in those EPDM–clay nanocomposites was also discussed based on the melt rheological data. The researchers concluded that enhanced dispersion of the silicate layers could be obtained in EPDM matrices with a higher Mooney viscosity (higher molecular weight) due to higher shearing during melt mixing. They also noted that linear polymer chains could be more intercalated into the interlayer spaces of organoclay than more highly branched EPDM chains. Zhen and coworkers studied the solid stage viscoelasticity of EPDM–clay nanocomposites with different types of organic modifiers. The addition of organoclay to EPDM results in a reduction of tan dmax (i.e., the tan d value at the glass transition temperature (Tg)) and a shift of Tg toward higher values. The decrease in tan d and a shift to higher Tg is most obvious for the exfoliated type EPDM–clay nanocomposite with methyl bis(2-hydroxyethyl)coco alkylamine modified montmorillonite (C12-Mt). Based on these results, they concluded that there was strong interaction between the EPDM and the surface of the silicate layers of C12-Mt. Chang and coworkers also reported a reduction of the tan d value and a shift of Tg in an EPDM–clay nanocomposite with liquid EPDM as a compatibilizer [2].

CHARACTERISTICS OF EPDM–CLAY NANOCOMPOSITES

475

15.5.3 Stability of EPDM–Clay Nanocomposites Dispersion or exfoliation of clay silicate layers is anticipated to provide flame retardation by reducing the mass of fuel, thermal insulation, and obstructing the release of combustible gases. Also, thermal degradation of polymer chains can be expected to be retarded by dispersion or exfoliation of clay silicate layers. Kang et al. studied the thermal stability of EPDM–clay nanocomposites by thermogravimetric analysis [8]. Weight loss (degradation) of pure EPDM began near 300 C and the curve declined gradually until approximately 400 C, when it decreased sharply. The addition of organoclay increased the initial weight loss (degradation) temperature over the pure EPDM compound. For different nanocomposite systems, the thermal stability for the composition with the maleic anhydride modified EPDM (EPDM-g-MA) as a compatibilizer was found to increase the greatest in comparison to simple EPDM–organoclay composites. They concluded that the inorganic silicate layers of the clay did not burn as readily and possibly interrupted the path of volatile degradation products and oxygen. The researchers then evaluated the flammability of the EPDM–clay nanocomposite using a conecalorimeter test. The peak heat release rate (PHRR) and the peak mass loss rate (PMLR) of the EPDM–organoclay nanocomposite with EPDM-g-MA were lower than those of the pure EPDM compound. The addition of the organoclay helped to reduce flammability because the dispersed silicate layers blocked the fuel source, which was generated by thermal decomposition of the EPDM polymer. Acharya and coworkers reported the effects of clay content on the thermal stability of EPDM–clay nanocomposites [11]. The initial decomposition temperature (Ti) for 2 wt.% loss in mass improved with increasing clay content. The Ti for neat EPDM was 325 C and that of the EPDM–clay nanocomposite with 2, 3, 4, 6, and 8 wt.% organoclay (hexadecyl amine modified montmorillonite) was 341, 355, 372, 372, and 374 C, respectively. Morlat-Therias et al. studied the photooxidation of EPDM–clay nanocomposites [7]. They found that the induction time of photooxidation was reduced in the presence of montmorillonite and the effect was enhanced in the case of the exfoliated nanocomposite. After the induction period, oxidation continued at a rate that was independent of the presence of montmorillonite. 15.5.4 Swelling Properties of EPDM–Clay Nanocomposites The swelling properties of EPDM–clay nanocomposites were investigated to study the interaction between EPDM and clay silicate layers. Such nanocomposites are expected to reduce swelling of rubber parts in the presence of oil in industrial applications. Kang et al. studied the swelling properties of EPDM–clay nanocomposites when exposed to toluene and also measured the solubility for those nanocomposites [8]. An EPDM–organoclay composite prepared without any organic modifier or compatibilizer and using a peroxide cure swelled up to about 300%, with the soluble matter determined to be about 5.3%. These data indicated a low level of effective cross-links. An EPDM–clay nanocomposite prepared using the maleic anhydride modified

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RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

EPDM (EPDM-g-MA) compatibilizer showed 200% swelling and had a soluble matter content of about 2.9%. The researchers concluded that these data indicated a higher level of effective cross-links, most likely due to interaction between the polymer chains and the filler. Therefore, EPDM nanocomposites containing EPDMg-MA absorbed less solvent. Acharya et al. studied the swelling properties of an EPDM–clay nanocomposite prepared with hexadecyl ammonium ion treated montmorillonite (16Me-Mt) [12]. The solvent (toluene) uptake ability of the EPDM–16Me-Mt nanocomposite not only increased with respect to neat EPDM but also initially increased with increasing addition of organoclay, reaching its highest level at 3 wt.% organoclay. They also measured the EPDM–16Me-Mt nanocomposite after equilibrium swelling in toluene using X-ray diffraction and compared the results to corresponding nanocomposites. These X-ray diffraction studies indicated that for nanocomposites containing more than 3 wt.% organoclay, when the solvent penetrated through the interlayer space, the domain space collapsed. They also showed that the amount of EPDM chains and the degree of exfoliation inside the clay layer were likely to be the same regardless of the total clay content. The researchers concluded that the silicate layers in the nanocomposites containing 3 wt.% organoclay provided a large free volume for toluene insertion. Further increases in organoclay content led to more ordered arrangements and decreased the availability of free volume, resulting in reduced solvent uptake. Zheng studied the cross-link density of EPDM–organoclay nanocomposites by gathering rapid solvent-swelling measurements (toluene for 72 h at 25 C) and applying the Flory–Rehner equation [16]. The cross-link densities of EPDM– organoclay nanocomposites were reduced by addition of organoclay. These results were in good agreement with the curing behavior. The vulcanization time (t90) of the nanocomposite was prolonged with increasing organoclay content, presumably due to the adsorption of curing agents on the filler surface. As expected, the decrease in cross-link density of the nanocomposite with increasing organoclay content was accompanied by an increase in mean molecular weight (Mc). 15.5.5 Mechanical Properties of EPDM–Clay Nanocomposites Tensile properties were evaluated by many research groups [1–4,6,8–11,13– 17,19,20]. Table 15.2 shows representative results for various EPDM–clay nanocomposites. Gatos et al. compared tensile properties of EPDM nanocomposites prepared on a two-roll mixing mill (open mill) to those of nanocomposites prepared in an internal mixer (internal mixer) [3]. The tensile strength and modulus (100%, 300%) of EPDM–clay nanocomposites prepared in an internal mixer showed higher values than those obtained for EPDM–clay nanocomposite prepared in an open mill. The higher shear rate and temperature of an internal mixer result in a greater dispersion state of the silicate layers in nanocomposites than can be achieved with an open mill. Ma et al. studied tensile properties of four types of EPDM–clay nanocomposites with different ethylene and diene content [6]. The elongation at break of

477

Ma et al. [6]

Gatos et al. [3]

References

1.1 1.0 1.0 1.0 1.0 1.0 1.4 1.8

– 1.8 – 1.6 1.7 1.4 2.8 2.9

1.5 1.6

4.9 7.1

EP33 EP33/organoclay J4045 J4045/organoclay J2080 J2080/organoclay J3080 J2080/organoclay

2.0

6.6

50% EPDM/50% EPDM-g-MA/ organoclay (open mill) EPDM/organoclay (internal mixer) 50% EPDM/50% EPDM-g-MA/ organoclay (internal mixer)

1.5

3.9

300% (MPa)

EPDM/organoclay (open mill)

System

100% (MPa)

Modulus

1.4 2.5 1.7 2.7 1.9 12.8 4.1 14.9

2.6 3.0

4.9

3.0

Tensile Strength (MPa)

TABLE 15.2 Tensile Properties of EPDM–Clay Nanocomposites and Related Samples

170 418 240 505 338 956 392 562

520 645

395

380

Elongation at Break (%)

EP33: Iodide value 26 Ethylene content 52% J4045: Iodide value 25 Ethylene content 52.5% J2080: Iodide value 12 Ethylene content 67% J2080: Iodide value 12 Ethylene content 67% (continued)

EPDM-g-MA:EPDM-grafted maleic anhydride

Comments

478 0.95 0.97 1.05 1.09 1.15

EPDM þ 2 wt.% EPDM þ 3 wt.% EPDM þ 4 wt.% EPDM þ 6 wt.% EPDM þ 8 wt.%

organoclay organoclay organoclay organoclay organoclay

0.92

EPDM, organoclay, and EPDM-g-MA EPDM and modified organoclay EPDM, EPDM-g-MA and modified organoclay

Pure EPDM

2.10 2.28 2.20

EPDM and organoclay

Kang et al. [8]

Acharya et al. [11]

2.07

System

100% (MPa)

– 1.73 2.24 2.35 2.43

–

3.23 4.23 3.98

3.20

300% (MPa)

Modulus

References

TABLE 15.2 (Continued)

2.4 2.9 3.4 3.4 3.5

1.6

8.92 7.45 8.55

8.90

Tensile Strength (MPa)

187 222 377 381 396

137

724 574 668

770

Elongation at Break (%)

Tensile strength and elongation at break values read from Acharya’s figure

Modified organoclay: a liquid-type, maleic anhydride-grafted vinyl polybutadiene treated montmorillonite

Comments

PREPARATION AND CHARACTERISTICS OF EPM–CLAY NANOCOMPOSITES

479

nanocomposites was improved relative to values obtained for pure EPDM. The addition of organoclay decreased the cross-link density, thus increasing macromolecule chain slippage during the stretching process. In addition, the elongation at break of EPDM nanocomposites which were prepared using EPDM with higher iodine values were lower than those of nanocomposites prepared from EPDM with lower iodine values. The researchers concluded that these results could also be attributed to the decrease in cross-link density. The improvement in the tensile strength of nanocomposites with high ethylene content was larger than that observed for nanocomposites with lower ethylene content. It was concluded that extension promoted the orientation of the silicate layers and induced crystallization of the polyethylene (PE) segment in the EPDM–clay nanocomposites with high ethylene content. Highly oriented amorphous chains resulting from strain-induced crystallization of the PE segments and the orientation of the clay layers in the nanocomposites with high ethylene content are likely responsible for the larger improvement in tensile strength when compared to the values obtained for nanocomposites with low ethylene content. Kang et al. studied tensile properties of EPDM nanocomposites, which were prepared using two types of compatibilizers (EPDM-g-MA, a maleic anhydridegrafted EPDM and LVPB-g-MA, a liquid type maleic anhydride-grafted EPDM) [8]. The EPDM nanocomposite prepared using the LVPB-g-MA modified organoclay resulted in good dispersion as determined by XRD analysis. This nanocomposite also showed higher modulus (100%, 300%) and reduced elongation at break values. The nanocomposite prepared using EPDM-g-MA also showed higher modulus and decreased elongation at break values compared to the EPDM and organoclay composite without compatibilizer. These data support the notion that an increased dispersion state of clay silicate layers is expected to result in nanocomposites with increased modulus and reduced elongation at break. Acharya et al. studied tensile properties of EPDM–clay nanocomposites with differing clay content [11]. Both tensile strength and elongation at break increased with increasing organoclay content up to 4 wt.%, and the corresponding improvement in either case was about 2.4 times with respect to that of the neat EPDM polymer. Upon further incorporation of organoclay in EPDM, both tensile strength and elongation at break remained nearly identical to the values obtained for the nanocomposite with 4 wt.% organoclay, indicating that the aggregation tendency of the clay silicate layers in the EPDM matrix gradually dominated.

15.6 PREPARATION AND CHARACTERISTICS OF EPM–CLAY NANOCOMPOSITES Hasegawa et al. prepared EPM–clay nanocomposites (EPM–CNs) by a melt intercalation method that contained different levels of octadecyl amine modified clay (C18-Mt) and studied the characteristics of these nanocomposites. The nanocomposites were prepared from maleic anhydride modified EPM (EPM–MA) polymers containing 0.42 wt.% maleic anhydride (4.81 mg KOH g  1) and with molecular weights of 125,000 and 397,000 as determined by GPC measurement.

480

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

FIGURE 15.4 Representative stress–strain curves: (a) EPM–CNs; (b) conventional composites.

15.6.1 Tensile Properties of EPM–CNs Representative stress–strain (S–S) curves of the EPM–CNs are shown in Figure 15.4. The S–S curve of EPM–MA shows typical behavior for an elastomer and has two yield points, the first around 80% strain and the second at the maximum just before the breaking point. In this discussion, the first yield point is defined as the point where the slope of the S–S curve decreases rapidly in less than 100% strain regions. At the first yield point, EPM polymer chains are thought to slip in the process of polymer chain rearrangement. Then the stress increases again by extension of the polymer chains between entanglement points in the larger strain region, and the EPM–MA is broken at around 900% strain after passing through the second yield point. The first yield points in the S–S curves of EPM–CNs are indistinct and become less pronounced with increasing clay content. Because rearrangement of the EPM polymer chain is thought to be inhibited by the silicate layers dispersed at the nanometer level, strain hardening of EPM–CNs begins in a smaller strain region compared to EPM–MA. The results of tensile tests are summarized in Table 15.3. TABLE 15.3 Tensile Properties of EPM–Clay Nanocomposites and Related Samples Sample

Filler

EPM–MA EPM–CN3 EPM–CN6 EPM–CN8 Composite 1 Composite 2 Composite 3 Composite 4 Composite 5

– C18-Mt C18-Mt C18-Mt Carbon black Carbon black Carbon black Talc Talc

Inorganic Modulus Strength at Elongation at Content (wt.%) (MPa) Maximum (MPa) Break (%) 0 2.9 6.1 8.3 5 15 30 4.9 10.1

7.2 11.2 17.2 23.2 7.9 12.8 18.0 7.1 7.5

4.4 3.7 4.7 4.2 4.3 4.2 4.8 3.6 3.4

900 550 400 130 860 600 400 820 720

PREPARATION AND CHARACTERISTICS OF EPM–CLAY NANOCOMPOSITES

481

The tensile moduli of EPM–CNs increase with increasing clay content. The tensile modulus of EPM–CN8, for example, is three times higher than that of EPM–MA. The elongation at break values for EPM–CNs, however, decrease with increasing clay content. It is thought that restriction of polymer chain rearrangement by dispersed silicate layers enhances the moduli but decreases the elongation at break of EPM–CNs. The tensile strength at the first yield point increases with the content of clay, as shown in Figure 15.4a, due to the increase of the moduli. But for the tensile strength at the second yield point, EPM–CN6 shows the highest value, while EPM–CN8 has a lower value because the elongation at break of EPM–CN8 is extremely low. The strength at the second yield point of EPM–CN3 is lower than that for EPM–MA because there is minimal restriction of polymer chain mobility by the silicate layers in EPM–CN3 and the elongation at break is smaller compared to EPM–MA. Figure 15.4b shows the S–S curves of conventional composites containing carbon black or talc in EPM–MA. The S–S curve of EPM–CN6 is also shown for comparison. EPM–CN6 exhibits much higher modulus and much smaller elongation than conventional composites at a similar filler loading level (5 wt.%). The S–S curve of EPM–CN6 is similar to that of the composite with 30 wt.% carbon black loading. In the case of conventional composites, the first yield points in the S–S curves also become less pronounced and the elongations at break decrease with increasing content of carbon black. 15.6.2 Temperature Dependence of Dynamic Storage Moduli of EPM–CNs The temperature dependence of the dynamic storage moduli of EPM–CNs shown in Figure 15.5 are higher than those of EPM–MA in the range from  150 to 100 C. The relative storage moduli of EPM–CNs and the conventional composites compared to those of EPM–MA are shown in Figure 15.6. The relative storage moduli of EPM–CNs below glass transition temperatures (Tgs, around  35 C as decided from

FIGURE 15.5 Storage moduli of EPM–CNs.

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RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

FIGURE 15.6 Relative storage moduli of EPM–CNs and conventional composites.

FIGURE 15.7 Storage moduli versus inorganic content: (a) 20 C; (b) 80 C.

the peak top of tan d) were relatively small. They drastically increased above Tgs, went through the peak tops around  10 C and then decreased. Figure 15.7 shows the relationship between the storage moduli and inorganic content of EPM–CNs and conventional composites at 20 and 80 C. The storage modulus of EPM–CN6 at 20 C was comparable to the values for conventional composites with 30 wt.% inorganic component loading (Figure 15.7a). These results indicate that the organoclay dispersed at the nanometer level exhibits a five times higher reinforcement effect compared to conventional fillers. EPM–CNs also showed an approximately five times higher reinforcement effect at 80 C (Figure 15.7b). 15.6.3 Creep Properties of EPM–CNs Creep properties were remarkably improved in EPM–CNs. Creep elongations of EPM–CNs were smaller compared to EPM–MA (Figure 15.8).

PREPARATION AND CHARACTERISTICS OF EPM–CLAY NANOCOMPOSITES

483

FIGURE 15.8 Creep test results for EPM–CNs and related samples.

The elongation of EPM–CN6 was restricted to less than 1% at 30 h, while EPM–MA was elongated over 50% at 1 h and then was broken within 2 h. The composite with 5 wt.% carbon black did not exhibit any noticeable enhancement in creep resistance; it was elongated and broken within 3 h. The composite with 30 wt.% carbon black showed enhancement in creep resistance nearly equal to EPM–CN6. Hasegawa et al. proposed that maleic anhydride groups grafted to the EPM polymer chains were selectively adsorbed on the dispersed silicate layers and formed strong ionic interactions due to the affinity of maleic anhydrides for ionic surfaces. They presumed that the silicate layers dispersed at the nanometer level would be bridged with some polymer chains containing maleic anhydride groups. Thus, the dispersed silicate layers act like large pseudo cross-link points and improve the creep resistance of EPM–CNs. 15.6.4 Swelling Properties of EPM–CNs Hasegawa and coworkers studied the swelling properties of EPM–CNs in hexadecane. Figure 15.9a–c shows the increments in weight, in length on the plane, and in thickness of the samples (10  10  2 mm) soaked in hexadecane at 25 C, respectively. The degree of swelling of EPM–CNs was drastically restricted compared to EPM–MA. The weight increments of EPM–CNs decreased with increasing clay content and were also much smaller. The weight increment of EPM–MA is over 1700%, while that of EPM–CN8 is restricted to 333%. The length increments on the plane in EPM–CNs also become remarkably smaller with increasing clay content. The length increment on the plane in EPM–CN8 is restricted to 45% compared to 170% for EPM–MA. On the other hand, the thickness increments are barely restricted by organoclay loading. Figure 15.10 shows the swelling increment (1500 h soak) as a

484

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

FIGURE 15.9 Swelling test results: (a) weight increments; (b) length increments (on the plane); (c) thickness increments.

function of inorganic content for both EPM–CNs and conventional composites. The restricting effect on swelling by organoclays is remarkably superior to that provided by conventional fillers, especially in weight and in length on the plane. For example, the weight increment of EPM–CN3 is approximately equal to that of the conventional composites with about a 20 wt.% filler loading.

PREPARATION AND CHARACTERISTICS OF EPM–CLAY NANOCOMPOSITES

485

FIGURE 15.10 Swelling increments versus inorganic content: (a) weight increments; (b) length increments (on the plane); (c) thickness increments.

486

RUBBER–CLAY NANOCOMPOSITES BASED ON OLEFINIC RUBBERS (EPM, EPDM)

FIGURE 15.11 Anisotropy between length and thickness increments versus inorganic content.

Figure 15.11 shows the anisotropy between the length increments on the plane and the thickness increments. EPM–CNs have a much higher anisotropy between the length increments on the plane and the thickness increments than those observed for the conventional composites. The ratio in EPM–CN8 is more than 2.5, for example, while the conventional composites with less than 10 wt.% filler loading do not exhibit any anisotropies. It is thought that not only the polymer chains but also the dispersed silicate layers at the nanometer level are orientated in parallel with a sheet plane prepared by compression molding. This orientation of both the silicate layers and the polymer chains is believed to selectively restrict the swelling increment in the length on the plane in EPM–CNs.

15.7 CONCLUSIONS Many researchers have succeeded in preparing olefinic rubber (EPDM, EPM)–clay nanocomposites, which are extremely important materials from the standpoint of industrial use, because olefinic rubber is the most widely used rubber in the world. These clay nanocomposites exhibit highly desirable properties such as enhanced mechanical and gas barrier properties, improved flame resistance and thermal stability, and reduced swelling and creep. Consequently, these nanocomposites have attracted worldwide attention and are the subject of research by several major chemical manufacturing companies.

REFERENCES 1. Usuki A.; Tsukigase, A.; Kato M. “Preparation and properties of EPDM clay hybrids”. Polymer, 43, 2185–2189 (2002).

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2. Chang, Y. W.; Yang, Y. C.; Ryu, S.; Nah, C. “Preparation and properties of EPDM organomontmorillonite hybrid nanocomposites”. Polym. Int., 51(4), 319–324 (2002). 3. Gatos, K. G.; Thomann, R.; Karger-Kocsis, J. “Characteristics of ethylene propylene diene monomer rubber/organoclay nanocomposites resulting from different processing conditions and formulations”. Polym. Int., 53(8), 1191–1197 (2004). 4. Gatos, K. G.; Apostolov, A. A.; Karger-Kocsis, J. “Compatibilizer effect of grafted glycidyl methacrylate on EPDM/organoclay nanocomposites”. Mater. Sci. Forum, 482, 347–350 (2005). 5. Hasegawa, N.; Okamoto, H.; Usuki, A. “Preparation and properties of ethylene propylene rubber (EPR)–clay nanocomposites based on maleic-anhydride modified EPR and organophilic clay”. J. Appl. Polym. Sci., 93(2), 758–764 (2004). 6. Ma, Y.; Wu, Y. P.; Wang, Y. Q.; Zhang, L. Q. “Structure and properties of organoclay/ EPDM nanocomposites: influence of ethylene content”. J. Appl. Polym. Sci., 99(3), 914–919 (2006). 7. Morlat-Therias, S.; Mailhot, B.; Gardette, J. L.; Da Silva, C.; Haidar, B.; Vidal, A. “Photooxidation of ethylene-propylene-diene/montmorillonite nanocomposites”. Polym. Degrad. Stab., 90(1), 78–85 (2005). 8. Kang, D. H.; Kim, D.; Yoon, S. H.; Kim, D.; Barry, C.; Mead, J. “Properties and dispersion of EPDM/modified-organoclay nanocomposites”. Macromol. Mater. Eng., 292(3), 329–338 (2007). 9. Mohammadpour, Y.; Katbab, A. A. “Effects of the ethylene-propylene-diene monomer microstructural parameters and interfacial compatibilizer upon the EPDM/montmorillonite nanocomposites microstructure: rheology/permeability correlation”. J. Appl. Polym. Sci., 106(6), 4209–4218 (2007). 10. Ma, Y.; Wu, Y. P.; Zhang, L. Q.; Li, Q. F. “The role of rubber characteristics in preparing rubber/clay nanocomposites by melt compounding”. J. Appl. Polym. Sci., 109(3), 1925–1934 (2008). 11. Acharya, H.; Pramanik, M.; Srivastava, S.K.; Bhowmick, A. K. “Synthesis and evaluation of high-performance ethylene-propylene-diene terpolymer/organoclay nanoscale composites”. J. Appl. Polym. Sci., 93(5), 2429–2436 (2004). 12. Acharya, H.; Srivastava, S.K. “Influence of nanodispersed organoclay on rheological and swelling properties of ethylene propylene diene terpolymer”. Macromol. Res., 14(2), 132–139 (2006). 13. Acharya, H.; Srivastava, S. K.; Bhowmick, A. K. “Ethylene propylene diene terpolymer/ ethylene vinyl acetate/layered silicate ternary nanocomposite by solution method”. Polym. Eng. Sci., 46(7), 837–843 (2006). 14. Acharya, H.; Kuila, T.; Srivastava, S. K.; Bhowmick, A. K. “Effect of layered silicate on EPDM/EVA blend nanocomposites: dynamic mechanical, thermal, and swelling properties”. Polym. Compos., 29(4), 443–450 (2008). 15. Zheng, H.; Zhang, Y.; Peng, Z. L.; Zhang, Y. X. “Preparation and properties of semitransparent EPDM/montmorillonite nanocomposites”. Polym. Polym. Compos., 13(1), 53–60 (2005). 16. Zheng, H.; Zhang, Y.; Peng, Z. L.; Zhang, Y. X. “Influence of clay modification on the structure and mechanical properties of EPDM/montmorillonite nanocomposites”. Polym. Test., 23(2), 217–223 (2004).

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17. Gatos, K. G.; Karger-Kocsis, J. “Effects of primary and quaternary amine intercalants on the organoclay dispersion in a sulfur-cured EPDM rubber”. Polymer, 46(9), 3069–3076 (2005). 18. Jeon, H. U.; Lee, D. H.; Choi, D. J.; Kim, M. S.; Kim, J. H.; Jeong, H. M. “Characteristics of rubber/sodium montmorillonite nanocomposites prepared by a novel method”. J. Macromol. Sci. Phys., 46(6), 1151–1163 (2007). 19. Wu, Y. P.; Ma, Y.; Wang, Y. Q.; Zhang, L. Q. “Effects of characteristics of rubber, mixing and vulcanization on the structure and properties of rubber/clay nanocomposites by melt blending”. Macromol. Mater. Eng., 289(10), 890–894 (2004). 20. Acharya, H.; Srivastava, S.K.; Bhowmick, A.K. “Synthesis of partially exfoliated EPDM/ LDH nanocomposites by solution intercalation: structural characterization and properties”. Compos. Sci. Technol., 67(13), 2807–2816 (2007). 21. Usuki, A.; Kato, M.; Okada, A.; Kurauchi, T. “Synthesis of polypropylene–clay hybrids”. J. Appl. Polym. Sci., 63(1), 137–138 (1997). 22. Kato, M.; Usuki, A.; Okada, A. “Synthesis of polypropylene oligomer—clay intercalation compounds”. J. Appl. Polym. Sci., 66(9), 1781–1785 (1997). 23. Kawasumi, M.; Hasegawa, N.; Kato, M.; Usuki, A.; Okada, A. “Preparation and mechanical properties of polypropylene  clay hybrids”. Macromolecules, 30(20), 6333–6338 (1997). 24. Varghese, S.; Karger-Kocsis, J.; Gatos, K. G. “Melt compounded epoxidized natural rubber/layered silicate nanocomposites: structure–property relationships”. Polymer, 44 (14), 3977–3983 (2003). 25. Nieuwenhuizen, P. J.; Van Duin, M.; Haasnoot, J. G.; Reedijk, J.; Mcgill, W. J. “The limiting value of ZDMC formation: new insight into the reaction of ZNO and TMTD”. J. Appl. Polym. Sci., 73(7), 1247–1257 (1999). 26. Nieuwenhuizen, P. J. “Zinc accelerator complexes: versatile homogeneous catalysts in sulfur vulcanization”. Appl. Catal. A Gen., 207(1), 55–68 (2001). 27. Dirksen, A.; Nieuwenhuizen, P. J.; Hoogenraad, M.; Haasnoot, J. G.; Reedijk, J. “New mechanism for the reaction of amines with zinc dithiocarbamates”. J. Appl. Polym. Sci., 79(6), 1074–1083 (2001). 28. Nielsen, L. E. “Models for the permeability of filled polymer systems”. J. Macromol. Sci. Chem., 1(5), 929–942 (1967). 29. Usuki, A.; Hasegawa, N.; Kato, M.; Kobayashi, S. Polymer–clay nanocomposites. In Advances in Polymer Science, Inorganic Polymeric Nanocomposites and Membranes, Springer-Verlag, Berlin, Heidelberg, 2005, pp. 150–151.

CHAPTER 16

RUBBER–CLAY NANOCOMPOSITES BASED ON THERMOPLASTIC ELASTOMERS JOSEPH H. KOO OFODIKE A. EZEKOYE JASON C. LEE WAI K. HO MORGAN C. BRUNS

16.1 INTRODUCTION Thermoplastic elastomer (TPE) is a family of rubber-like materials that can be processed and recycled like thermoplastic materials, unlike conventional vulcanized rubbers, which cannot be recycled and requires cross-linking. Thermoplastic polyurethane (TPU), one type of TPE is becoming increasingly important as an engineering material because it displays properties of both elastomers and thermoplastics. In particular, TPUs have numerous applications such as coatings, adhesives, foams, rubbers, and thermoplastic elastomers [1]. Conventional TPU, however, is known to exhibit poor resistance to heat. This poor performance in heat resistance limits its applications [2]. The introduction of inorganic nanomaterials as additives into polymer systems has resulted in polymer nanostructured materials exhibiting a multiplicity of high performance characteristics beyond what traditional polymer composites possess [3,4]. These improved properties include higher thermal and flame resistance, moisture and chemical resistance, decreased permeability, charge dissipation, and thermal/electrical conductivity. The objective of this chapter is to thermally characterize thermoplastic polyurethane elastomer nanocomposites (TPUNs) at different heating conditions. The effects of weight loadings of layered clays on the thermal and flammability performance of this novel class of materials are explored using a variety of test protocols and methods such as thermogravimetric

Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

489

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RUBBER–CLAY NANOCOMPOSITES BASED ON THERMOPLASTIC ELASTOMERS

analysis (TGA), radiant panel experimental apparatus, UL 94, and cone calorimeter. The TPUNs are intended to develop for propulsion system application. Research on thermoplastic polyurethane–clay nanocomposites is ubiquitous; however, most of the studies focus on the enhancement of mechanical and thermophysical properties of the nanocomposites. There are only limited studies that investigate the thermal and flammability properties of TPU–clay nanocomposite. Tortora et al. [5] synthesized polyurethane–clay nanocomposites using diphenylmethane diisocyanate, poly(e-caprolactone), di(ethylene glycol), and poly (e-caprolactone) organically modified montmorillonite. They found that there was an improvement in the elastic modulus and yield stress but a decrease in the stress and strain at breaking takes place on increasing the clay content. Tien and Wei [6] improved the morphology of the layered silicates–polyurethane nanocomposites by using reactive swelling-agent-modified silicates as pseudochain extenders for polyurethane prepolymer. It was observed that there is a 34% increase in Young’s modulus, a 1.7-fold increase in the tensile strength, and a 1.3-fold increase in the elongation to break in the TPUN containing only 1 wt.% trihydroxyl group swellingagent-modified silicates. In another study, Tien and Wei [7] found a 44
C increase in glass transition temperature of the hard segmented phase and a 2.8-fold increase in the storage modulus. They also enhanced the heat resistance and degradation kinetics. A 40
C increase in the degradation temperature and a 14% increase in the degradation activation energy occurred in polyurethane containing 1 wt.% trihydroxyl swelling agent-modified montmorillonite compared to that of the pristine polyurethane. Berta et al. [8] investigated the effect of polyol molecular weight and functionality on nanodispersion of clay in PU–clay nanocomposites and their thermal and combustion properties. They found that the dispersion of organoclays in polyurethane nanocomposites tends to improve with molecular weight of the polyol whereas it seems unaffected by increasing its functionality from 2 to 2.5. The nanocomposites volatilization is also delayed in air owing to the barrier effect of the clay layers. They reported that the formation of char during combustion lowers the peak heat release rate in cone calorimetry and eliminates fire-induced dripping of the nanocomposite sample during the UL 94 test. Marchant et al. [9] conducted a study of flame retardant thermoplastic polyurethane nanocomposites with montmorillonite nanoclays, carbon nanofibers, and polyhedral oligomeric silsesquioxane (POSS ). They found that the coefficient of thermal expansion of nanoclay TPUN increased by twofold for 10 wt.% nanoclay. There is a 71% reduction of peak heat release rate for TPU–5% nanoclay nanocomposite as compared with the pristine TPU when exposed to a heat flux of 50 kW/m2 using cone calorimetry. The thermal conductivity of the TPU–5% nanoclay nanocomposite slightly increased relative to the neat sample from 0.202 to 0.239 W/(m
C) measured at 45
C. Groups at the U.S. National Institute of Standards and Technology (NIST) and Cornell University both reported that nanocomposites alone, with no other flame retardant, reduced the parent polymer’s flammability and enhance char formation. Giannelis [10,11] found self-extinguishing properties of polycaprolactone (PCL)

SELECTION OF MATERIALS

491

and aliphatic-polyetherimide layered silicate nanocomposites when the nanocomposites were exposed to small open flame tests. Researchers at NIST used cone calorimetry and radiative gasification to show that polyamide-6 (PA6), polystyrene (PS), and maleated polypropylene (PP-g-MA) montmorillonite (MMT) nanocomposites had enhanced char formation and gave up to 75% lower flammability, as measured by the peak heat release rate (PHRR) or peak mass loss rate [12–14]. In most cases, the carbonaceous char yield was limited to 2–5% mass fraction; consequently, the total heat release rate (THRR) was not affected significantly. In a recent review, Gilman [15] concluded that a definite flame retardant effect appears in the form of a reduction in PHRR of polymers that contain nanodispersed clay. This is generally true for most thermoplastic polymers studied to date: PS, PA6, PP, polyethylene (PE), ethylene-co-vinyl acetate (EVA), polyamide-12 (PA12), and poly(methyl methacrylate) (PMMA); however, it appears that different mechanisms are operative in different polymers. Depending on the polymer, the clay may change the decomposition products; it may cause cross-linking and ultimately catalyze carbonaceous char formation. In some cases, no char is formed and it is only the quality of the clay char that controls the flame retardant effect. Gilman further concluded that it is necessary to combine the nanocomposite with other flame retardant additives to pass regulatory flammability tests when ignition properties dominate. Morgan [16] has recently reviewed a combination of conventional flame retardants with polymer–clay nanocomposites that yielded success in various regulatory flammability tests. In this chapter, two types of commercially available thermoplastic polyurethane elastomer were melt-blended with various loadings (2.5–10 wt.%) of montmorillonite nanoclay using twin-screw extrusion. The morphological, physical, thermal, flammability, thermophysical properties, and kinetic parameters of these two families of polymer–clay nanocomposites were characterized. The processing– structure–property relationships of this class of novel thermoplastic polyurethane elastomer–clay nanocomposites were established.

16.2 SELECTION OF MATERIALS 16.2.1 Polymer Resin The properties of TPU are strongly correlated to the relative amounts of hard and soft segments in this segmented block copolymer. This two-phase structure is generally specified to occur in so-called hard and soft domains in which the strong polar hard segments are separated from the nonpolar soft segments. The hard domains are often thought of as nanometer length scale fillers within the often-amorphous soft domains. The relative proportions of hard and soft segments impact the overall crystalline nature of these domains. It is not surprising then that nanofiller modification of TPU will induce interesting mechanical and thermal modifications. Two commercially available TPUs were used in this study. Initially, Pellethane 2102-90A, a polyester polycaprolactone-based TPU elastomer [17] was selected for

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investigation. This is a TPU material manufactured by Dow Plastics and has the following properties: outstanding abrasion resistance; good low temperature flexibility and impact resistance; good resistance to fuels, oils, and most nonpolar solvents; good hydrolytic stability; range of hardness; transparency; low compression set; high compressive strength; low gas/vapor permeability; relatively high moisture vapor transmission rates; easy processibility with conventional extrusion and thermoplastic molding techniques; and low extractable levels. Sequentially, a new TPU, Desmopan DP 6065A [18] manufactured by Bayer MaterialScience was selected. Desmopan DP 6065A is an aromatic polyether-based TPU. It can be processed by injection molding. It is characterized by good hydrolysis and microbe resistant, short cycle times, and is plasticizer free. Typical applications include rigid–flexible composite systems and sport shoe soles. Material properties of Pellethane 2102-90A and Desmopan DP 6065A are listed in Table 16.1 [17,18].

TABLE 16.1 Material Properties of Pellethane 2102-90A and Desmopan DP 6065A Typical Properties, (US Units, SI Metric Units) General Specific gravity Hardness (Shore A) Taber abrasion, H-18 wheel, 1000 g load, 1000 cycles, mg loss Bayshore resilience (%) Mold shrinkage at 100 mil thickness: flow direction/cross-flow direction (in./in., mm/mm) Mechanical Tensile strength (psi, MPa) Tensile stress at 100% elongation (psi, MPa) Tensile stress at 300% elongation (psi, MPa) Ultimate elongation (%) Tear strength (Die C) (psi, kN/m) Flexural modulus at 23
C/at  30
C (psi, MPa) Compression set, % (22 h at 70
C/22 h at 23
C) Thermal Vicat softening temperature (Rate A) (
F,
C) Glass transition temperature (DMA) (
F,
C)

Pellethane

Desmopan

1.2 94A 10

1.09 65A 76

0.006/0.007

49 0.008/0.008

6500, 44.8 1600, 11.0

1790, 12.3 380, 2.6

4000, 27.6

610, 4.2

440 830, 145 12000/82.7

890 290, 50.9 1440, 9.9/2400, 16.6

25, 29

34, 12

243, 117

133, 56

 15,  26

 49,  45

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Desmopan DP 6065A is much softer than the previously tested Pellethane 2102-90A TPU. Both materials are compounded and injection molded into 10.2- by 10.2- by 0.32-cm test specimens, they are transparent materials. A type A durometer is used to test the shore hardness on two stacked 0.36 cm (1/8 in.) thick panels as specified under ASTM-2240. The measured Pellethane 2102-90A shore hardness, 95A, closely matches the one that is listed in the material data sheet, 94A. The measured shore hardness of Desmopan that ranges from 67A to 70A is bit higher than the listed hardness, 65A. The material at the center of the panel was found to be softer than that at the edges (more discussion of material hardness will be presented later in the chapter). 16.2.2 Nanoparticles Southern Clay Products montmorillonite nanoclays were used in the work reported. These nanoparticles will reinforce the polymer in the nanoscale and will enhance the thermal stability and mechanical properties of the polymer nanocomposites (PNCs) as shown by Pinnavaia and Beall [19], Koo [3], and Bhattacharya et al. [20]. Insertion of a polymer into nanoclays by exfoliation has been shown to increase mechanical properties, barrier performance, and application processing [19]. To achieve exfoliation, MMT clays are surface treated to be compatible with the polymer and high shear mixing efficiency of the dispersing apparatus. Exchanging MMT clay inorganic counterions, sodium, with long chain quaternary ammonium ions results in surface treatment of MMT clay with satisfactory hydrophobic characteristics and ease of dispersing the surface-treated nanoclays within the hydrophobic polymer continuous phase [3,19,20]. For this work, TPUs samples are loaded with 2.5, 5, 7.5, and 10 wt.% of Cloisite 30B MMT nanoclays. Cloisite 30B is natural MMT clay modified with a quaternary ammonium salt manufactured by Southern Clay Products. Specific chemical pretreatment modifier (MT2EtOH, methyl tallow bis-2-hydroxyethyl quaternary ammonium) was used for the preparation of Cloisite 30B [21,22].

16.3 EXPERIMENTAL 16.3.1 Processing of Thermoplastic Elastomer Nanocomposites A 30 mm Werner Pfleider corotating twin-screw extruder, whose L/D can be varied from 21 to 48 with options for multiple feeds and vents, was used to melt-blend about 9.2 kg (20 lbs) of each formulation. The TPU resin pellets, which were desiccant dried before compounding, were fed via the main feed throat. While the nanoclay was fed into the twin-screw extruder downstream, the nanoclay was injected into the molten resin via the main feed throat in order to maximize dispersion given the 90 s residence time. All nanoclays were injected as received from the vendor. The screw speed was 200 rpm and residence time was about 90 s. The processing temperature was approximately 20
C above its melting temperature, immediately before it was

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TABLE 16.2 Material Matrix for Baseline and TPU Nanocomposites Formulation 1 2 3 4 5 6 7 8 9 10 11

Polymer Matrix

Filler

EPDM (88%) Pellethane 2102-90A (100%) Pellethane 2102-90A (97.5%) Pellethane 2102-90A (95%) Pellethane 2102-90A (92.5%) Pellethane 2102-90A (90%) Desmopan DP 6065A (100%) Desmopan DP 6065A (97.5%) Desmopan DP 6065A (95%) Desmopan DP 6065A (92.5%) Desmopan DP 6065A (90%)

12% Kevlar None 2.5% Cloisite 30B nanoclay 5% Cloisite 30B nanoclay 7.5% Cloisite 30B nanoclay 10% Cloisite 30B nanoclay None 2.5% Cloisite 30B nanoclay 5% Cloisite 30B nanoclay 7.5% Cloisite 30B nanoclay 10% Cloisite 30B nanoclay

injection molded to produce the test specimens. In the Phase I study, Pellethane 210290A TPU was selected and in the Phase II study, Desmopan DP 6065A TPU was selected. Different loading levels of Cloisite 30B were melt-blended in both TPUs. The baseline Kevlar -filled EPDM rubber, pure TPU, and TPUN formulations used in this research program are listed in Table 16.2. 16.3.2 Morphological Characterization Luo and Koo [23] developed a general method to quantify the dispersion of nanoparticles in polymer matrices using transmission electron microscopy (TEM) analyses. The dispersion quantity, D, is defined as the probability of inclusion particle free-path spacing falling into a certain range of the mean spacing m, according to the particle spacing data frequency distribution. Two quantities, D0.1 and D0.2 are proposed, which are the probabilities of the particle free-path spacing falling into the ranges of m  0.1 m and m  0.2 m, respectively. Both normal and lognormal distributions are discussed and in both cases, the quantities D0.1 and D0.2 are specified as monotonically increasing functions of m/s, where m and s are the mean particle free-path spacing and standard deviation, respectively. The dispersion parameter, D0.1, based on the measurement of the free-path spacing distance between the single clay sheets [24], carbon nanofibers and multiwall carbon nanotubes [25] from the TEM images are presented by Luo and Koo. This methodology was adapted in this study. In past work, selected Pellethane TPU blends with the baseline Kevlar-filled EPDM were tested for ablation performance using the Air Force Research Laboratory (AFRL) Pi-K char motor with aluminized propellants for ablation. The Pellethane–5% Cloisite 30B TPUN was ranked highest in terms of ablation resistance and mechanical properties for their respective family of nanoparticles, thus chosen for more detailed thermophysical and kinetic properties characterization in this study. TEM analyses, the Pellethane–5% Cloisite 30B TPUN was also analyzed using wide-angle X-ray diffraction (WAXD) for dispersion, no peaks were observed in WAXD. The WAXD and TEM analyses based on the Luo and Koo method [23,24]

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confirmed that the Pellethane–5% Cloisite 30B TPUN achieved exfoliation in the polymer matrix. 16.3.3 Thermal Properties Characterization 16.3.3.1 Experimental Method The required data for calculating the kinetic parameters by TGA are temperature, weight loss, and the rate of weight loss. These data are obtained with a Perkin Elmer Model TGA7 thermogravimetric analyzer. It has a vertical design with a high sensitivity microbalance and electric platinum wound furnace, which is capable of heating rates up to 180
C/min to a maximum temperature of 1000
C. The microbalance is located above the furnace and is thermally isolated from it. A hang-down wire is suspended from the balance down into the furnace. At end of the hang-down wire is the platinum sample pan. The sample temperature is measured with a chromel–alumel thermocouple located inside the furnace. The purging gas flow rate is regulated by a mass flow controller and is kept at 20 mL/min. 16.3.3.2 Calibration Because the thermocouple and sample pan are physically separated, the indicated and actual sample temperatures differ. Temperature calibration as well as baseline run is required each time when the heating rate is changed. The temperature calibration is performed according to the Curie point of magnetic materials. Curie point is the temperature at which a ferromagnetic material becomes paramagnetic; due to thermal energy, the dipoles cannot line up as they can in the ferromagnetic state. Four reference materials such as alumel, nickel, perkalloy, and iron are used to perform a calibration across the usable temperature range (100–900
C). The reference sample and sample pan are tared. A magnet is placed beneath the sample. Below the reference Curie point, the mass balance measures the magnetic force. As the material is heated past the reference Curie point the force on the mass balance decreases. If the temperature correction at any Curie point is more than 2
C, the limits on the temperature control unit is changed, and this Curie calibration is repeated. Furnace calibration is also performed in accordance to the user manual. A baseline run will be the same procedure as a sample run except with an empty sample pan. 16.3.4 Flammability Properties Characterization There is amazing similarity between how a sample ablates and its flammability properties. One measure of the ablative performance of a material is its ability to absorb heat when thermally assaulted. Often, a scenario-dependent parameter, the heat of ablation, is used to relate the incident heat flux to the mass flux of material leaving the ablator. A well-designed ablator will have a mass loss rate that is well tuned to the overall sample mass and to the heating environment. In flammability analysis, the ASTM E1354 (cone calorimeter) is used to measure the heat release rate (HRR) of materials. Recognizing that the heat release rate is in fact proportional to the mass loss rate of volatiles liberated from the sample, one sees that inferences of the relative ability of samples to liberate volatiles can be made using cone calorimetry.

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16.3.4.1 Radiant Panel Test TPUN samples of 10.2 by 10.2 by 0.32 cm (4 by 4 by 1/8 in.) are tested using a radiant panel. The samples are tested at low heat flux between 35 and 90 kW/m2 and the front and backside temperatures are measured [26,27]. These experiments are able to provide detailed thermal data on the behavior of the TPUN materials. 16.3.4.2 UL 94 Test A standard flammability test performed on polymers is the UL 94 test [28]. Of the three types of UL 94 test setups: surface burn, vertical burn, and horizontal burn tests; the vertical burn test was run on the TPUN formulations. A 2.54 cm (1 in.) flame is placed at 45
to the bottom of a clamped 12.7  1.27  0.32 cm (5  0.5  0.125 in.) test specimen. After 10 s the flame is removed and the time the specimen remains on fire is recorded. A second 10 s flame is introduced and the time of burn after removal is again recorded. The specimen is held 30.48 cm (12 in.) above a piece of cotton. The cotton is used to determine the flammability of any drip from the material. Based on these observations the specimen can either pass or fail V0, V1, or V2. V0 is the most stringent of the three. The specimen must not burn for more than 10 s and is not allowed to drip and ignite the cotton. V1 allows up to a 30 s burn but still restricts drip that ignites the cotton. V2 also allows less than 30 s burn time, however, allows drip that ignites the cotton. 16.3.4.3 Cone Calorimeter Cone calorimetry is a standard test method for characterizing heat and visible smoke release rates for materials and products using oxygen consumption calorimetry [29]. The operating principle is that, in most materials, the heat release rate is directly proportional to the mass rate of oxygen consumed during burning. After ignition by a spark ignition source, test specimens burn in ambient air conditions subject to a predetermined external heat flux (typically 0–100 kW/m2; 50 kW/m2 in this study). Test results are expressed in terms of time to ignition; peak, average and total heat release rates; mass loss and mass loss rate; effective heat of combustion; visible smoke development; and release rates of carbon monoxide and carbon dioxide. In this work, samples were examined with no additional conditioning/drying prior to testing. Measurements were made on a FTT Dual Cone Calorimeter with an exhaust flow of 24 L/s using the standardized cone calorimeter procedure (ASTM E-1354-10) [29]. All samples were tested without frame and grid with each sample’s backside wrapped in aluminum foil. Data collected from all samples have an error of 10% based on a specimen surface area of 100 cm2. Three repetitions were conducted for each blend and specific attention paid to the heat release rate and peak heat release rate. 16.3.5 Thermophysical Properties Characterization Thermophysical properties of TPU and TPUNs were evaluated. Density measurement was performed using a pycnometer. Specific heat capacity measurement was performed using a differential scanning calorimeter (DSC). Thermal diffusivity measurement was performed using laser flash method. Thermal conductivity was

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calculated using data from density, specific heat capacity, and thermal diffusivity measurements. 16.3.5.1 Density True volume of the material was measured using a pycnometer. By measuring the weight of the sample and employing Archimedes’s principle of gas displacement and the technique of gas expansion, the density of the sample was obtained. 16.3.5.2 Specific Heat Differential scanning calorimetry (DSC) was performed to obtain the specific heat of the TPU and TPUNs. A small amount of sample/material sealed in a pan was heated at a set heating rate in a nitrogen environment, and an analyzer recorded the heat flow needed to keep the temperature of the sample pan increasing at the same rate as the reference or empty pan. Away from conditions in which kinetic decomposition occurs, the specific heat was calculated by dividing the heat flow by the sample weight and the heating rate. 16.3.5.3 Thermal Diffusivity Laser flash diffusivity at 45
C and 55
C was performed to obtain the thermal diffusivity of the TPU and TPUNs. A laser is used to supply a high-intensity, short-duration pulse of thermal energy to one face of a sample. The intensity of the beam was controlled by varying the laser power supply and use of attenuating filters. The resulting temperature rise of the other face of the sample was monitored using an infrared detector. Then, the thermal diffusivity was determined from a numerical analysis of the infrared detector output.

16.4 NUMERICAL 16.4.1 Modeling of Decomposition Kinetics 16.4.1.1 Modeling Approaches There are many different ways that one can model the thermal degradation of polymeric materials. The modeling approaches vary in their predictive abilities and in the number of calibration parameters required to fit any given type of model. Clearly, the type of polymeric material also, in general, affects the modeling approach that can be undertaken. Within the past two decades, a push has been made to model polymeric thermal degradation at smaller resolved scales and closer to the more fundamental quantum chemical pathways. Because of the time and length scales that are typically observed in experimental polymeric degradation processes (1000s of seconds and millimeter length samples), it is clearly infeasible to simulate polymeric degradation at the quantum-mechanical time and length scales. Nevertheless, important information is being gained using modeling approaches such as reactive molecular dynamics that couple to ab initio calculations. In many reactive molecular dynamics simulations, classical molecular dynamics codes are modified to include potentials that are able to transition between predictions of classically continuously bonded atoms to reactive modifications derived from quantum chemistry calculations. Examples in the thermoplastic degradation include

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the work of Nyden and coworkers [30,31]. Such work is progressing toward being able to describe some aspects of nanomodification to thermoplastic degradation. At even higher (coarser) levels of description are population balance based descriptions of polymer degradation in which a polymeric system is described in terms of a population of elements that carry specific descriptors of the polymer system characteristics. At the simplest level, this could be a description of the molecular weight distribution of the polymeric system. In more advanced cases, the polymeric system might be formulated into a set of different chemical families for which descriptors of conformation, molecular weight, and so on might be used to parameterize transition and kinetic rules of change between elements of the population. Population-based models of polymer degradation still require some information about the fundamental kinetic rates which might either come from reactive molecular dynamics simulations or from extrapolation of kinetic rates made from gas phase systems [32]. Linear polymer degradation has a somewhat simple and elegant mathematical structure for analysis in the population balance modeling framework. Branched polymer systems represent somewhat greater mathematical structure. In these systems, the overall connectivity of the branched system must be modeled. Graph-based mathematical models and other mathematical structures relevant to network systems (e.g., percolation modeling) have been used to simulate the loss of connection associated with these systems. An example of work on network degradation associated with branched/thermoset degradation is work by Galina and Lechowicz [33–35]. Similar to linear polymer degradation processes, there is a need for fundamental kinetic rates to identify the rate of breakage of the bonds in the networks. 16.4.1.2 Fitting Thermoplastic Polyurethane Decomposition Kinetics As previously noted, TPU are composed of soft and hard domains primarily comprised of soft and hard segments. The hard segments are held together by thermally modifiable hydrogen bonds, while the urethane backbone has strong covalent bonds. Several studies have investigated the degradation of TPU with the goal of identifying the likely chemical steps involved. Kang et al. show that the apparent activation energy for degradation in TGA experiments performed at 20
C/min increases with the percentage of hard segment content [36]. Thus, one would expect that if nanofillers act as hard segments, then a similar effect should take place by the inclusion of nanofillers. They find that the decomposition of TPU is a two-step process in which CO2 is released in the first step with H2O and HCN released in the second step. Rosu et al. perform simultaneous gas analysis as well as thermogravimetric and DSC analysis on TPU [37]. They indicate that a four-step degradation process is most likely the degradation path. They point out the overall complexity of the processes, but note that the likely first step in the breakage processes is the polyurethane scission and the formation of isocyanates and glycols. It would appear then that the decomposition steps identified by Kang et al. [36] represent an averaged or reduced/simplified form of the steps noted by Rosu et al. [37]. Given the overall complexity of identifying meaningful pathways for TPU degradation, it has been useful in polymer degradation work to look at model form fitting of kinetics.

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Standard methods to fit the overall decomposition processes of complex polymeric materials include differential methods, isoconversion methods, and optimization-based approaches. In a sense, all of these methods are in fact optimization based. The increasing power of computer codes has made it relatively easy to propose general functional forms for hypothesized degradation pathways and to use modern optimization techniques such as genetic algorithms, simulated annealing, and so on. to fit the parameters within the decomposition reaction [32,38]. Much of the error in the use of fitting strategies has been associated with improper fitting approaches [39]. In general, the rate of decomposition of a material might be modeled by the kinetic rate equation, 
n
1 dw w E 2 ¼ Af exp 2 ð16:1Þ w0 dt w0 RT where A is the preexponential factor, E is the activation energy, n is the order of reaction, and f(w/w0) is an arbitrary function of instantaneous weight fraction. Note that there are issues associated with decoupling intrinsic kinetic rates from a global mass loss rate equation because the effect of mass transfer could confound the analysis. Essentially, Eq. (16.1) indicates that mass loss and transport are significantly faster than the kinetic decomposition rates such that kinetics are rate limiting. To predict the thermal response of a material, accurate values of the kinetic parameters in Eq. (16.1), over the range of decomposition temperatures, are required in a thermal model. In the literature, there are several methods to determine the kinetic parameters from experimental data. One of the most widely used methods is multiple heating rate technique, proposed by Friedman [40] in 1965, through the use of TGA. TGA is a common polymer characterization technique to determine the thermal stability of a sample material by recording the weight change of the material as a function of increasing temperature or time at a constant rate. A furnace heats the sample while a sensitive microbalance monitors loss or gain of sample weight due to chemical reactions, decomposition, solvent or water evaporation, and oxidation. Normally, the measurement is carried out in air or in an inert environment, such as nitrogen or argon, at a constant purging rate. The multiple heating rate technique will provide the preexponential factor, activation energy, and order of reaction over a wide range of heating rates. Henderson [41] proposed a method for calculating the kinetic parameters from the Arrhenius kinetic rate equation. Taking the natural logarithm of both sides will result "
# 1 dw w E ln ¼ ln Af ð16:2Þ w0 dt w0 RT When plotting ln[(1/w0)(dw/dt)] as a function of 1/T for all heating rates, linear equations may be fitted for each parametric value of w/w0. These linear equations will have slopes of  E/R, and each equation will have y-intercept of ln[Af(w/w0)] for each parametric value of w/w0. Then, an average activation

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energy, Eavg., from all of the linear curve fits can be calculated and is used to obtain ln[Af(w/w0)] for each parametric value of w/w0. Calculate an average ln [Af(w/w0)] for each parametric value of w/w0 for all heating rates. Next, the arbitrary function of weight fraction can be defined as,
   w ðw2wf Þ n f ð16:3Þ ¼ w0 w0 where wf is the final weight of the material at the end of a reaction in the TGA run. Multiplying both sides by A and taking the natural logarithm will yield, 
    w ðw2wf Þ ln Af ¼ lnðAÞ þ n ln w0 w0

ð16:4Þ

When plotting ln[Af(w/w0)] as a function of ln[(wwf)/w0], linear equation(s) may be used to fit and obtain both A and n. Another method that is widely used to describe kinetics is the isoconversion method [42]. Since the mechanisms of material thermal degradation are often unknown and too complicated to be characterized by a simple kinetic model. Processes can occur in multiple steps having different rates. The basic idea of the method can be summarized as follows. The rate of degradation is a function of temperature and solid conversion, Xs, dXs ¼ f ðT; Xs Þ ¼ kðAs 2 Xs Þn dt

ð16:5Þ

where Xs is defined as Xs ¼

w0 2 w w0 2 wf

ð16:6Þ

As is the maximum weight loss attainable at each temperature (value has been taken as 1.0), and k is the kinetic constant, which corresponds to the Arrhenius equation
 E k ¼ A exp 2 ð16:7Þ RT A single-order reaction is used (n ¼ 1), and the Ozawa–Flynn–Wall method is considered, for obtaining the value of activation energy [43] ln b ¼ ln


 AE E 2 5:332 lnð12 Xs Þ2 1:05 R RT

ð16:8Þ

where b is heating rate. Plotting ln b as a function of 1/T, linear equations may be used to fit for each parametric value of Xs to obtain the values of activation energy and preexponential factor.

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These kinetic parameters imply that the thermal decomposition of the material can be expressed by a simple kinetic equation, with single effective values of A and E. By adjusting the order of reaction, a separate set of kinetic parameters can be obtained, but when plotting log A as function of E, a linear curve fit can be observed using different values of n, which is result of the kinetic compensation effect. We use standard fitting approaches to define the overall kinetic characteristics of polymer degradation. Note that the effective kinetics derived from TGA includes not only the polymer degradation kinetics but also include the rate of volatile mass loss by transport. The general assumption is that the mass loss rate is significantly faster than the polymer degradation kinetics such that the effective time is dominated by the degradation kinetics. As the heating rate increases, different kinetic processes may be affected and the mass loss rate may be of comparable time scales as the heating rate that would complicate interpretations of kinetics at the lower heating rates. In a later section, we discuss the applicability of extrapolating low heating rate TGA experiments to practical conditions.

16.5 DISCUSSION OF RESULTS 16.5.1 Nanoparticle Dispersion Polymer nanocomposites containing Cloisite 30B are classified under three categories: unmixed if the nanoclay platelets are packed tightly together, intercalated if the platelet stacks are open but still grouped, and exfoliated if each nanoclay platelets appears independent of the each other. All the Pellethane–clay nanocomposites were analyzed using WAXD for dispersion as shown in Figure 16.1. No peaks were observed in WAXD, which indicated that all loadings of nanoclays were exfoliated in 5000

Counts per second (CPS)

4500

Pellethane–2.5% Cloisite 30B Pellethane–5.0% Cloisite 30B

4000

Pellethane–7.5% Cloisite 30B Pellethane–10.0% Cloisite 30B

3500 3000 2500 2000 1500 1000 500 0 0

2

6

4

8

10

2θ

FIGURE 16.1 Wide-angle X-ray diffraction of Pellethane–clay TPUNs.

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FIGURE 16.2 TEM micrographs of Pellethane with (a) 2.5%, (b) 5%, and (c) 10% Cloisite 30B nanoclay TPUNs with unit bar of 100 nm.

the TPU polymer. TEM analyses were also performed on all TPUNs. The TEM micrographs, in Figure 16.2 showing magnification of 100 nm scale bar. They show single clay platelets, which were visible in the TPU polymer and confirmed that the 2.5, 5, and 10 wt.% Cloisite 30B nanoclay achieved intercalation/exfoliation in the TPU polymer matrix. Based on the Lou and Koo method [23] a morphological characterization was conducted on Pellethane TPU–clay nanocomposites. Table 16.3 shows the quantitative measurements of TPUN–clay in clay loadings for 2.5, 5.0, 7.5, and 10 wt.% [24]. The quantitative analyses show the D0.1 values of the 2.5, 5.0, 7.5, and 10.0 wt.% materials are 8.4, 8.6, 9.8, and 7.5%, respectively. Only the 10.0 wt.% TPUN–clay formulation is classified in intercalated state and the other TPUN–clay formulations are classified in exfoliated state. The pretest Desmopan–clay TPUN specimens change from transparent (neat Desmopan) into darker color as the percentage of nanoclay increases [44]. The WAXD plots of the four different nanoclay loading of Desmopan TPUNs are shown in Figure 16.3. No peaks were observed in WAXD, which indicated the all loadings of nanoclays were exfoliated in the TPU polymer. Certain material properties TABLE 16.3 Results of Quantitative Measurements of TPU and TPUN-Clay Materials

Polymer Matrix 2102-90A

Loading Level (wt.%)

Mean Free-Path Spacing, m (nm)

Standard Deviation, s (nm)

m/s

D0.1 (%)

2.5 5.0 7.5 10.0

105.1 66.3 45.0 27.8

110.6 68.0 39.9 33.4

0.9503 0.9750 1.1278 0.8323

8.4 8.6 9.8 7.5

Classification Exfoliated Exfoliated Exfoliated Intercalated

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500,000 Desmopan–2.5% Cloisite 30B

450,000 Counts per second (CPS)

Desmopan–5.0% Cloisite 30B

400,000 Desmopan–7.5% Cloisite 30B

350,000

Desmopan–10.0% Cloisite 30B

300,000 250,000 200,000 150,000 100,000 50,000 0 0

2

6

4

8

10

2θ

FIGURE 16.3 Wide-angle X-ray diffraction of Desmopan–clay TPUNs.

enhancement will not fully be achieved if the nanoparticles are not well dispersed within the polymer. Examination of this is done using transmission electron microscopy. No stacks are observed in any of the three 40 kX images in Figure 16.4, scale bar is 100 nm. The lack of platelet stacks indicates that the nanoclay is exfoliated and well dispersed in the Desmopan TPU material. 16.5.2 Thermal Properties Fundamental material property enhancements are important to assess a material’s viability for specific applications. In addition to flammability and mechanical strength, two very important material properties are density and hardness. Weight considerations are important in aerospace and technology applications. In addition, many manufacturing processes and applications depend heavily on the stiffness of a material.

FIGURE 16.4 TEM micrographs at 40 kX of Desmopan with (a) 2.5%, (b) 5%, and (c) 10% Cloisite 30B nanoclay TPUNs with unit bar of 100 nm.

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In processing Pellethane and Desmopan polymer nanocomposites, high weight loadings of different nanofillers have been shown to change the viscosity to such a degree that injection molding becomes difficult. Processing also uses weight percentages rather than volume percentages to determine the fraction of resin to nanoclay formulates. Densities of Pellethane 2102-90A and Desmopan DP 6065A materials with 2.5, 5.0, 7.5, and 10.0 wt.% of Cloisite 30B were measured using a pycnometer. The neat Pellethane TPU has a density of 1.18 g/cc and the density of the Pellethane–clay TPUN increases with the increase of Cloisite 30B loadings. Densities of 0, 5.0, 7.5, and 10.0 wt.% of Pellethane–clay TPU are 1.18, 1.2, 1.22, and 1.23 g/cc, respectively. The neat Desmopan TPU has a lower density of 1.1 g/cc as compared to Pellethane TPU. The Desmopan TPU’s density increases only slightly when 5 wt.% of Cloisite 30B is added (1.12 g/cc). However, there is a significant increase in density from 7.5 to 10 wt.% loadings (1.13 and 1.17 g/cc) [44]. Shore hardness tests were performed on Pellethane 2102-90A and Desmopan DP 6065A materials with 2.5, 5.0, 7.5, and 10.0 wt.% of Cloisite 30B using ASTM D2240 standard. In general, the hardness of Pellethane–clay TPUN increases slightly as the clay loading increases from 5 to 10 wt.% of Cloisite 30B. The shore hardness of neat Pellethane is 90A and Pellethane–10 wt.% clay is 95A. For the Desmopan DP 6065A samples, two sets of material sizes were used. One test used a stack of two materials with a thickness of 0.32 cm (1/4 in.) and the second test used a material that had been compression pressed and made to be 0.64 cm (1/2 in.) thick. For each of these materials, five different locations on the 10.2 by 10.2 cm (4 by 4 in.) top surface were measured and averaged. The hardness measured using the stacked materials appear stiffer. This is expected since the 0.64 cm (1/2 in.) thick materials are found to have porous air pockets at the interfaces of separate materials as they are heat compression pressed. The shore hardness of neat Desmopan is 60A and Desmopan–10 wt.% clay is 85A [44]. The Desmopan Cloisite 30B samples are tested at 10, 20, 40, and 100
C/min in nitrogen atmosphere. Samples were made by cutting the lengths of pellets to samples between 13 and 16 mg. Tests show that TGA results are sensitive to specimen size and shape. By using pellets of the same diameter and cutting them to the specified weight, the shape and size of the samples were kept consistent. Each formulation was tested at the four heating rates from 100 to 900
C. The results from the 20
C/min (CPM) heating rate tests are presented as representative of the other three heating rates, since the trend of each are similar [44]. Figure 16.5 shows the weight percentage versus temperature data of the neat Desmopan as well as the four Cloisite 30B formulations. The neat Desmopan degrades rapidly at a lower temperature than the Cloisite 30B formulations. This indicates the addition of Cloisite 30B that enhances the thermal stability of the polymer. Thermal stability increases as more nanoclay is added to the Desmopan. The 2.5% Cloisite 30B formulation clearly is much more stable than the neat Desmopan, and the 5% Cloisite 30B formulation is much more stable than the 2.5% Cloisite 30B formulation. However the 5, 7.5, and 10% Cloisite 30B profiles appear to be very close. Similar results were observed in the UL 94 flame spread test where the neat material melted but did not burn, the 2.5% Cloisite 30B material

DISCUSSION OF RESULTS

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Desmopan® with 30B at 20ºC/min Weight percent remaining

120 100 80 60 40 20 0 350

300

400

500

450

550

600

Temperature (ºC) 5.0% 30B

2.5% 30B

Neat

7.5% 30B

10% 30B

FIGURE 16.5 Thermogravimetric analysis of Desmopan–Cloisite 30B formulations at 20 CPM.

burned but fell apart, and the 5, 7.5, and 10% Cloisite 30B materials burned for roughly the same amount of time but did not fall apart (discussed later). The onset, 10% weight loss, and 50% weight loss temperatures shown in Figure 16.6, also show the trend that there is significant increase in thermal stability up to the 5 wt.% formulation [44]. Formulations with a higher weight loading are found to have roughly the same onset and 50% weight loss temperatures. The 10% weight loss temperature is found to be constant throughout all formulations; this means that the enhancements of the Cloisite 30B are not noticed until a higher Desmopan® with 30B at 20ºC/min

Temperature (ºC)

415 405 395 385 375 365 355 345 335 2

0

4

6

8

10

12

Weight percent loading Onset

10% weight lost

50% weight lost

Poly. (onset)

Linear (10% weight lost)

Poly. (50% weight lost)

FIGURE 16.6 Onset and weight loss temperatures of Desmopan–Cloisite 30B formulations at 20 CPM.

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506

temperature is reached. Since the Cloisite 30B degrades at a higher temperature than the Desmopan, as the Desmopan degrades, a larger percentage of the remaining material is Cloisite 30B. The nanoclay begins to build a larger protective layer preventing more of the Desmopan from decomposing. The rate of weight loss data show that one reaction is taking place during this experiment [44]. This result is somewhat different from what was noted by Kang et al. [36] and Rosu et al. [37]. Increasing the nanoclay loading decreases the peak rate of weight loss as well as increases the temperature at which these peaks occur. This can be observed in the TGA graph which shows that the increase in nanoclay loading stretches the area of largest degradation rather than simply shifting the graph. Addition of nanoclay additives have been shown to delay thermal degradation as well as help to create a char layer. Different heating rate experiments [44] are performed on both neat TPU and 5% Cloisite 30B TPUN from 20 to 500
C/min (Figure 16.7). There are two degradation reactions that occur. The distinct reactions are represented as the two discrete slopes observed in both materials at the lower heating rates. As heating rate increases the curves shift to the right. The shape of these curves is maintained in the case of the 5% Cloisite 30B TPUN, however, the two distinct slopes in the neat Desmopan TPU curves mold into one slope at the higher heating rates [44]. In this inert nitrogen environment, the thermal stability is most enhanced by the Cloisite 30B at the lower heating rates. The 50 wt.% decomposition temperature difference at the 20
C/min experiments is large, 37
C. As the heating rate increases this temperature difference decreases. In all TPUN experiments, the remaining 120 Neat Desmopan 20 CPM Neat Desmopan 50 CPM Neat Desmopan 100 CPM Neat Desmopan 250 CPM Neat Desmopan 500 CPM Desmopan–5% Cloisite 30B 20 CPM Desmopan–5% Cloisite 30B 50 CPM Desmopan–5% Cloisite 30B 100 CPM Desmopan–5% Cloisite 30B 250 CPM Desmopan–5% Cloisite 30B 500 CPM

Remaining weight (%)

100

80

60

40

20

0 200

250

300

350

500 450 400 Temperature (ºC)

550

600

650

700

FIGURE 16.7 Thermogravimetric analysis of neat Desmopan TPU and Desmopan–5% Cloisite 30B TPUN.

DISCUSSION OF RESULTS

507

weight percentage is roughly the same as the amount of Cloisite 30B that is added, and a char is created. 16.5.3 Flammability Properties Three types of flammability experiments will be discussed in this section: (a) IR radiant panel testing of a set of Pellethane–clay TPUNs using 35, 50, and 90 kW/m2 radiant heat fluxes, (b) UL 94 testing of a set of Desmopan–clay TPUNs, and (c) cone calorimeter testing of neat Pellethane TPU, neat Desmopan TPU, Pellethane–5% Cloisite 30B TPUN, and Demospan–5% Cloisite 30B TPUN at a radiant heat flux of 50 kW/m2. 16.5.3.1 IR Radiant Panel Test All Pellethane–clay TPUN samples are first exposed to a heat flux of 50 kW/m2, which represents a medium fire environment (simulating a large trash container fire) using the IR radiant panel test apparatus [26,27]. The surface temperatures are recorded using an IR pyrometer and backside temperatures are recorded by K-type thermocouple. Then the best TPUN– clay sample is exposed to additional heat fluxes of 35 and 90 kW/m2. Surface temperatures are only measured for heat flux of 50 kW/m2 test case. Effect of Nanoclay Weight Loading on Backside Temperature of TPUN The backside temperatures of Pellethane–clay TPUN samples with different weight loadings of Cloisite 30B [26,27] are shown in Figure 16.8. The samples were exposed to a 50 kW/m2 heat flux. All composite samples formed a char layer 450 Neat Pellethane

400

2.5% Cloisite 30B 5% Cloisite 30B

350 Temperature (ºc)

7.5% Cloisite 30B 10% Cloisite 30B

300 250 200 150 100 50 0 0

25

50

75

100 125 150 175 200 225 250

275 300 325

Time (s)

FIGURE 16.8 Backside temperatures of Pellethane–clay TPUNs at 50 kW/m2.

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RUBBER–CLAY NANOCOMPOSITES BASED ON THERMOPLASTIC ELASTOMERS

after the surface of the sample was burnt. This can be explained by the change in slope of the 5, 7.5, and 10% curves. When the material was exposed to the specified heat flux the temperature increases steadily for about 120 s. A char layer is then formed. This char layer acts as a thermal barrier and protects the polymer from the heat source and as a result the slope of temperature–time curve decreases. However, for the pristine Pellethane TPU (baseline) and 2.5% Cloisite 30B samples, the material melted and sloughed off the sample holder. The thermocouple on the back of the sample was exposed directly to the heat source, thus resulting in substantial increase in temperature. For the baseline sample, the polymer melted around 120 s, and the 2.5% Cloisite 30B sample melted around 130 s. The Pellethane–5% Cloisite 30B TPUN gives the lowest backside temperature or the best thermal insulation. All of these results can be explained by the interaction between the pristine Pellethane TPU and the nanometer clay layers. When the clay content increases, the clay may begin to aggregate and thus becomes a less effective thermal insulator. Effect of Heat Flux on Backside Temperature of TPUN–Clay Pellethane–5% Cloisite 30B TPUN exhibits the best insulation at 50 kW/m2. Therefore, it was exposed to additional heat fluxes of 30 and 90 kW/m2 to explore the effect of heat flux on the backside temperature [26,27]. Higher heat fluxes give higher backside temperatures. From 0 to around 150 s, there is not much difference in backside temperature between the different heat fluxes. However, after 150 s, the backside temperature of the sample under 90 kW/m2 increases dramatically. There is not much difference in backside temperature between 35 and 50 kW/m2. Surface Temperature of TPUN–Clay at 50 kW/m2 Averaged surface temperatures of Pellethane–clay TPUNs (three samples) with different weight loadings of Cloisite 30B were measured using an IR pyrometer [26,27]. The samples were exposed to a heat flux of 50 kW/m2. As discussed above, the 2.5% 30B formulation melted and sloughed off the stand; therefore, the data for 2.5% 30B formulation is not reported. It should be noted that the surface temperature of the samples increases significantly after the char layer is formed. Moreover, the surface temperatures did not increase any further after the char layer was fully formed. Therefore, the average surface temperature of the char layer is calculated from the formation till the end of the experiment. The average surface temperature difference between the 5, 7.5, and 10% weight loadings was small (10
C). This observation suggests that the char layer effectively protects the Pellethane–clay TPUN from the heat regardless of the weight loadings of the nanoclays. 16.5.3.2 Cone Calorimeter Test Cone calorimeter experiments at 50 kW/m2 are tested on the neat Desmopan TPU, neat Pellethane TPU, Desmopan–5% Cloisite 30B TPUN, and Pellethane–5% Cloisite 30B TPUN. Two or three experiments are performed to test for repeatability. The neat material experiments are the least repeatable due to the fact that the material melted and dripped off the test fixture, for this reason the test with the higher PHRR is used in the comparisons.

DISCUSSION OF RESULTS

509

2500 Neat Desmopan Desmopan–5% Cloisite 30B Neat Pellethane

HRR (kW/m2)

2000

Pellethane–5% Cloisite 30B

1500

1000

500

0 0

50

100

150

200

250

300

350

400

450

Time (s)

FIGURE 16.9 Heat release rate of Desmopan and Pellethane TPU and Cloisite 30B TPUNs.

The HRR graph of the materials is shown in Figure 16.9. The flammability properties are shown in Table 16.5. The material’s time to sustained ignition is not changed drastically with the addition of Cloisite 30B, þ 2 s for Pellethane TPUN and 1 s for the Desmopan TPUN. However, a dramatic decrease in PHRRs is observed. Pellethane TPU PHRR decreases by 73% and Desmopan TPU PHRR decreases by 50% with the addition of 5 wt.% Cloisite 30B. In the first 60 s, the HRR is higher in the TPUN. This is due to a higher rate of decomposition at this stage observed in the mass loss graph. This is consistent with the mass loss curves from the TGA experiments discussed in the previous section. The average HRR after 180 s is higher in the TPUs. 16.5.3.3 UL 94 Test Vertical UL 94 experiments are performed by hanging a 12.7 cm (5 in.) long specimen above a piece of cotton. The bottom portion of the specimen is exposed to a 2.54 cm (1 in.) flame for 10 s. After which the time the material continues to burn is measured. After the flame extinguishes the 2.54 cm (1 in.) flame is reintroduced for an additional 10 s. Again the time it takes for the flame to extinguish is measured. The total time of burn is used to determine the UL 94 rating of the material. In addition to burn time, whether or not the material drips and if the drip burns the cotton below then it is used to determine the material’s flammability rating. Upward flame spread can be characterized in a relatively simple form if the heat release rate of the sample is known [45]. Not surprisingly, as the heat release rate determined from cone calorimetry is reduced, the upward flame spread rate decreases. This clarifies the importance of reducing the heat release rate in order for a sample to be able to pass the UL 94 test.

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RUBBER–CLAY NANOCOMPOSITES BASED ON THERMOPLASTIC ELASTOMERS

TABLE 16.4 Summary of Vertical UL 94 Experiments

2.5 wt.% Cloisite 30B 5 wt.% Cloisite 30B 7.5 wt.% Cloisite 30B 10 wt.% Cloisite 30B

No. of Drops

1st Drop Time (s)

1st Drop Length (cm)

Last Drop Time (s)

Burn Time (s)

4–8 1 0 0

39 76 – –

1.3–3.8 10.2–11.7 – –

85 – – –

185 177 228 237

The neat Desmopan TPU material dripped quickly in this experiment [44]. In fact, the material dripped so quickly that the bulk material did not burn for more than 1 s. The melt that did drip was flammable, indicating that although the bulk material did not burn due to the melts low viscosity. If the bulk material was held in a horizontal position it would burn. This is confirmed when the Desmopan–Cloisite 30B TPUN samples are tested. The Cloisite 30B increases the viscosity of the TPUN and thermally shields the material. Even at low weight loadings, 2.5%, the burning material does not flow as it does in the neat sample. However, at this low weight loading large cracks appear along the surface. Cracks propagate from the edges toward the center. From sample to sample, at this loading, 4–8 specimens would drop. Higher weight loadings of Cloisite 30B result in a stronger burnt material. This is evident in the fact that the Desmopan–5% Cloisite 30B sample only has a single material drop. At 7.5 and 10 wt.% loading, no specimens dropped. The material specimen length and times of drop are shown in Table 16.4. The Desmopan–Cloisite 30B TPUNs have a tan pigment. While burning the material, the cracks propagate toward the center. The black charred surface contrasts the tan soft melt at the center. This shows that the surface chars and the center of the material melts, degasses, and allows the bulk material to burn. Higher weight loading chars are strong enough to hold up the material as the weak soft melt degasses. The posttest material char of the Desmopan–10% Cloisite 30B TPUNs are clearly noticeable. The remaining tan pigment at the top of each specimen is due to the fact TABLE 16.5 Summary of Cone Calorimetry Data at Irradiance Heat Flux of 50 kW/m2

Material Neat Pellethane Pellethane–5% Cloisite 30B Neat Desmopan Desmopan–5% Cloisite 30B

tig (s) 32 34 28 27

PHRR (kW/m2) 2290 664 (71% reduction) 1031 518 (50% reduction)

Avg. HRR, 60 s (kW/m2)

Avg. HRR, 180 s (kW/m2)

Avg. Eff. Hc (MJ/kg)

Avg. SEA (m2/kg)

406 560

653 562

30 25

237 303

228 442

515 376

27 28

311 256

tig ¼ time to sustained ignition; PHHR ¼ peak heat release rate; Avg. HRR ¼ average heat release rate after ignition; Avg. Eff. Hc ¼ effective heat of combustion; Avg. SEA ¼ average specific extinction area.

DISCUSSION OF RESULTS

511

that the materials are held by a clamp. As the material burns, the clamp acts as an insulator, such that the conduction through the center of the material is not significant enough to melt or char this segment. 16.5.4 Microstructures of Posttest Specimens 16.5.4.1 Radiant Panel Posttest Specimens Scanning electron microscopy (SEM) micrographs showing surface view are taken on the char layer for Pelletane–5% 30B clay TPUN after all samples were exposed to different heat fluxes for 300 s. SEM micrographs in low magnification (2 kX) of neat Pellethane TPU at heat flux of 35 (Figure 16.10a) and Pellethane–5% 30B TPUN at heat fluxes of 35 (Figure 16.10b), 50 (Figure 16.10c), and 90 kW/m2 (Figure 16.10d). It can be observed that there are more “ridges” in the surface structure of the Pellethane–5% 30B TPUN specimens in 2 kX magnification as the heat flux increases. At higher heat flux, the polymer matrix is burnt off, and only the clay galleries remain. These galleries constitute the char layer on the surface and act as ceramic heat shields to effectively delay heat penetration. The mechanisms of polymer–clay nanocomposites were well documented by Gilman [15].

FIGURE 16.10 SEM micrographs at 2 kX showing surface view of Pellethane (a) pure TPU exposed to 50 kW/m2, (b) with 5 wt.% Cloisite 30B exposed to 35 kW/m2, (c) with 5 wt.% Cloisite 30B exposed to 50 kW/m2, and (d) with 5% Cloisite 30B exposed to 90 kW/m2.

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RUBBER–CLAY NANOCOMPOSITES BASED ON THERMOPLASTIC ELASTOMERS

FIGURE 16.11 SEM micrographs of Desmopan–10% Cloisite 30B at three burn stages: (a) heated surface, (b) 30 s burn surface, and (c) complete burn surface.

The TPUN–clay material enhances the thermal properties by forming a char layer on the surface and delays the heat penetration. The char layer is constituted of clay layers for the TPUN–clay material. The thermal conductivity properties of the TPUN–clay materials also contribute to the overall effectiveness of the TPUN–clay material as an insulator. 16.5.4.2 UL 94 Posttest Specimens As described in the previous section, flammability processes of failure appear to be multistepped. Initially the material begins to crack due to thermal stresses. As the material begins to degas, bubblelike structures are visible with the naked eye changing the surface roughness of the material. The existence of these bubbles suggests that for this scale of mass loss process, the actual mass loss model might be important. SEM images of the surface of Desmopan–10% Cloisite 30B at different stages of a UL 94 experiment are shown in Figure 16.11. The heated surface, Figure 16.11a, is a surface that has not burned, however, is near a flame. The heat from the flame causes the material to begin to crack. A surface that begins to burn for 30 s and externally extinguished is shown in Figure 16.11b. Large degassing bubbles are observed. The material continues to burn, until no combustible material remains more bubbles appear (Figure 16.11c). 16.5.5 Thermophysical Properties Thermophysical properties including density, specific heat, thermal diffusivity, thermal conductivity, coefficient of thermal expansion of pure Pellethane TPU, and Pellethane–5% Cloisite 30B TPUN were characterized [46,47]. Thermal conductivity was derived based on the relationship: k ¼ a(T)r(T)cp(T), where k is the thermal conductivity, a is the thermal diffusivity, r is the density, and cp is the specific heat. All thermophysical properties data were tabulated in Table 16.6. As expected, the 5% Cloisite 30B TPUN (TPUN–clay) has higher density than the neat TPU. The specific heat of the Pellethane–clay TPUN is substantially higher than the neat Pellethane TPU. The thermal conductivity value of the Pellethane–clay TPUN is higher than the neat Pellethane TPU, and is more thermally conductive than the neat Pellethane TPU. The coefficient of thermal expansion (CTE) of the pure

DISCUSSION OF RESULTS

513

TABLE 16.6 Thermophysical Properties of Kevlar-Filled EPDM, TPU, and TPUN

Material Kevlar-filled EPDMa Neat Pellethane TPUb 5% Cloisite 30B TPUNb a b

Density (g/cc)

Specific Heat (J/(g
C))

Thermal Diffusivity (10  7 m2/s)

Thermal Conductivity (W/(m
C))

1.16 1.18 1.21

1.423 1.514/1.524 1.814/1.886

2.196 1.133/1.007 1.087/1.057

0.363 0.202/0.180 0.239/0.241

Kevlar-filled EPDM properties were measured at room temperature. TPU/TPUNs properties were measured at 45
C/55
C.

Pellethane TPU and Pellethane–clay TPUN of various loadings are measured. The CTE of Pellethane–clay TPUN increases as the Cloisite 30B loading increases. All these thermophysical properties need to be characterized in the charred state at elevated temperatures. Using relatively simple ideas it would not be apparent that the changes in the thermophysical properties would positively impact the thermal, flammability, and ablative performance of the samples. 16.5.6 Kinetic Parameters Thermogravimetric data of the materials are obtained at six different heating rates of 5, 10, 20, 50, 100, and 150
C/min [46,47]. Three formulations are tested in a nitrogen atmosphere: Kevlar-filled EPDM, pure Pellethane TPU, and Pellethane–5% Cloisite 30B TPUN. Weight loss, rate of weight loss, and kinetic parameters using the isoconversion method for all three formulations are presented below. 16.5.6.1 Kevlar-Filled EPDM The fraction of weight remaining (weight loss) and rate of weight loss versus temperature for Kevlar-filled EPDM are calculated for six different heating rates [46,47]. Fraction of weight remaining is plotted until 13% remained mass because no additional mass is lost as the temperature increases. Observed from the curve of weight loss, the onset temperature of decomposition shifts to higher temperatures with increasing heating rates. At the lowest heating rate of 5
C/min the onset temperature is approximately 430
C while at the highest heating rate 150
C/min it is approximately 500
C. This is not a surprising result for this type of reaction. One notes that when the reaction rate is in fact of Arrhenius form, then a Damkohler number (characteristic heating time divided by kinetic time) decreases with increasing heat rate. The effect of this is a relative slowing of the overall reaction progress, which results in a delay in the temperature and time at which the primary decomposition reaction takes place. At low heating rate (5
C/min), the peak weight loss rate occurs at a lower temperature, at approximately 484
C (757 K), and if the heating rate is 30 times larger (150
C/min), the peak weight loss rate occurs at approximately 612
C (885 K), a 16.9% increase. Thermal stability provides information about the degradation resistance of a material to thermal load. The activation energy, calculated from

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RUBBER–CLAY NANOCOMPOSITES BASED ON THERMOPLASTIC ELASTOMERS

180 Kevlar-filled EPDM Pure TPU TPU 5% Cloisite 30B

Activation energy (kJ/mol)

160 140 120 100 80 60 40 20 0 0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1

Solid conversion (Xs)

FIGURE 16.12 Activation energy as a function of solid conversion of Kevlar-filled EPDM, pure TPU, and TPU 5% Cloisite 30B calculated by isoconversion method.

isoconversion method, as a function of solid conversion of Kevlar-filled EPDM is depicted in Figure 16.12. Until solid conversion of about 0.13, activation energy values are increasing with decreasing solid conversion, dictating the slow weight loss regime. When pyrolysis occurs, the activation energy reaches a peak value and slowly decays up to conversion of 0.90, and then it decreases further as it reaches complete conversion. The apparent average activation energy is 141.416 kJ/mol, and the apparent preexponential factor is calculated to be 4.50  1013 min1. 16.5.6.2 Pure Pellethane TPU The fraction of weight remaining and rate of weight loss versus temperature for pure Pellethane TPU are calculated for six different heating rates [46,47]. The pyrolysis reaction initiates at approximately 300
C (573 K). At low heating rate (5
C/min), the peak weight loss rate occurs at a lower temperature, at approximately 343
C (616 K), and if the heating rate is 30 times larger (150
C/min), the peak weight loss rate occurs at approximately 461
C (734 K), a 19.1% increase. The calculated activation energy by isoconversion method as a function of solid conversion of pure Pellethane TPU is depicted in Figure 16.12. Unlike Kevlar-filled EPDM, pure Pellethane TPU degrades at lower temperatures. Activation energy value peaks at conversion of 0.04 at approximately 108 kJ/mol and decreases slowly. The apparent average activation energy is 96.747.9 kJ/mol, and the apparent preexponential factor is calculated to be 4.29  107 min1. The activation energy found in this study is consistent with the work of Rosu et al. [37] who show the apparent activation energy as a function of conversion progress and also in terms of their more detailed four-step reaction path. They find

DISCUSSION OF RESULTS

515

that the activation energy is between 104 and 125 kJ/mol for the first half of the decomposition process. The preexponential value is comparable to the Rosu et al. values at a ¼ 0.2. Both our data and the Rosu et al. data are quite different from the data of Kang et al. [36] in which the apparent activation energy is in the range of 20–50 kJ/mol. 16.5.6.3 Pellethane–5% Cloisite 30B TPUN The fraction of weight remaining and rate of weight loss versus temperature for Pellethane–5% Cloisite 30B TPUN are calculated for six different heating rates [46,47]. From the weight loss rate curves, two reactions are observed. Fully exfoliated nanoclay particles increase the decomposition temperature as well as the temperature of peak weight loss rate. At low heating rate (5
C/min), the peak weight loss rate occurs at approximately 410
C (683 K), a 10.9% increase in terms of thermal stability compare to pure Pellethane TPU. When the heating rate is 30 times larger (150
C/min), the peak weight loss rate occurs at approximately 544
C (817 K), a 19.6% increase by heating rate and an 11.3% increase at the same heating rate compare to pure Pellethane TPU. The activation energy calculated by isoconversion method as a function of solid conversion of Pellethane–5% Cloisite 30B TPUN is depicted in Figure 16.12. Unlike Kevlar-filled EPDM and pure Pellethane TPU, the activation energy follows two regimes because two reactions occur. The activation energy required to complete the first reaction has an average value of 99.9 kJ/mol. From solid conversion of 0.37 onward, the activation energy steadily increases and peaks near complete conversion. Over this second region, the apparent average activation energy is 104.57.7 kJ/mol, and the apparent preexponential factor is calculated to be 3.80  108 min  1. Based on temperatures at peak weight loss rates, Kevlar-filled EPDM outranked the proposed formulations at all heating rates. This observation contradicts with the SRM ablation experiments performed earlier. However, the percentage of temperature increases of peak weight loss rate from low heating rate (5
C/min) to high heating rate (150
C/min), or the thermal stability from 5 to 150
C/min, showed that the proposed formulations were better more stable than the neat formulation (19.6% for Pellethane–5% Cloisite 30B TPUN but only 16.9% for Kevlar-filled EPDM). This trend falls in line with our SRM ablation data which makes us wonder if it is valid to use low heating rate data for applications involving high heating rate. Further studies are in progress to address this issue. The validity of extrapolation of heating rates in TGA experiments to high heating conditions is frequently questioned [46,47]. An extreme condition is the atmospheric entry scenario where polymeric ablators are used as part of the thermal protection system. We can estimate the heating rate in a typical ablator by noting that the maximum temperature difference, DTMAX, in a steady-state ablation process is on the order of 3000
C, the recession rate, s_ , is approximately 0.1 mm/s, and the thickness of the pyrolysis zone can be estimated as d  a=_s. Thus, the scaling for the temperature rise rate is s_ 2 DTMAX =a. For a nominal case, for which the thermal diffusivity (a) is 0.1  10  6 m2/s, we see that the temperature rise rate is 300
C/s (18,000
C/min) which is approximately 120 times larger than our fastest heating rates. The rate of heating is a practical limitation on use of TGA for ablation property characterization.

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Using similar scaling ideas, we discuss the validity of the isothermal specimen assumption for TGA samples as a function of heating rate. Assuming a sample characteristic length scale of d, which is of order 1 mm for TGA samples, the time scale for diffusion over that length is t DIFF  d 2 =a. The TGA heating rate in degrees per second is b. The product of the heating rate and characteristic diffusion time (bd2/a) represents the characteristic temperature difference across the sample. The highest heating rate that we use is 150
C/min (2.5
C/s). The characteristic temperature difference across the sample is thus approximately 25
C. While this seems to be a large temperature difference, recognize that typical temperatures for degradation are approximately 400
C (673 K). As such, the relative temperature error, based on absolute temperatures, is less than 4% at the highest heating rates that was used. Perhaps of greater consequence for high heating rate conditions is the impact of the mass loss processes on the effective degradation kinetics. As previously noted, the effective kinetics derived from TGA includes not only the polymer degradation kinetics but also include the rate of volatile mass loss by transport. Exploring such issues is fertile ground for future investigations.

16.6 SUMMARY AND CONCLUSIONS Formulations of Pellethane and Desmopan TPUs with the addition of various loading (2.5–10 wt.%) of Cloisite 30B nanoclay were blended via twin-screw extrusion. WAXD, TEM, and dispersion analyses of TEM micrographs confirmed that the nanoclays were fully exfoliated in the 2.5, 5, and 7.5 wt.% of TPUN–clay samples and intercalated in the 10 wt.% TPUN–clay sample. The Pellethane–5 wt.% Cloisite 30B TPUN formulation ranked well when ablation and mechanical properties were considered, and were thus chosen to be the candidates to replace the current state-ofthe-art Kevlar-filled EPDM for the next generation of solid rocket motor internal insulation material [44,46–54]. Dynamic thermogravimetric analyses in nitrogen at six different heating rates were performed to obtain the kinetic parameters for Pellethane–clay TPUN as well as the baseline Kevlar-filled EPDM and pure Pellethane TPU. Weight loss and weight loss rate measurements were recorded. All formulations exhibited expected behavior. As the heating rate was increased, the temperature of weight loss shifted to higher values, delaying the solid conversion. Closed-loop verification showed that both the modified Friedman and isoconversion methods provided reasonably accurate kinetic parameters for the benchmarking materials, but in the calculations for the proposed Pellethane–clay TPUN formulations, the modified Friedman method sometimes produced negative activation energy and order of reaction that was clearly unrealistic. The isoconversion method did not encounter such problem, and it was adapted for the kinetics calculations in this study. Values of the apparent activation energy and preexponential factor were obtained and are generally consistent with values found in the open literature. Physical, thermophysical, and flammability properties characterization were performed including density, hardness, specific heat, thermal diffusivity, thermal conductivity, coefficient of thermal expansion, heat release rate,

NOMENCLATURE

517

smoke concentration, and gas generation. The low peak heat release rate, smoke concentration, and CO emission also made the TPUN–clay a very attractive class of fire retardant polymers that has potential to be used for commercial markets. These material properties characterization will provide more insight on the thermal response of the TPU material. It appears that the TPUN–clay material enhances the thermal properties by forming a ceramic-like char layer on the surface and delays thermal penetration. The char layer is constituted of clay layers for the TPUN–clay material. Also, it has been shown in other work (e.g., Herrera-Alonso et al.) that nanoclay addition to TPU decreases gas permeability and increases the Tg of the material. The permeability  Tg reinforcement is assumed to derive from the clay strengthening of the soft segments of the PU. Continuation of this work should include the proper method of precharring and tailoring the material for specific heat and thermal diffusivity measurements [51,52]. Once the thermophysical properties are obtained, they should be extrapolated for even higher temperatures. TGA experiments at high heating rates (up to 500
C/min) on several of the TPUNs are reported to investigate the effects of heating rate and nanoparticle loading [44,46,47]. Advanced mechanistic models of thermoplastic degradation using optimization [32] with the TGA data will be employed to determine kinetic parameters.

16.7 NOMENCLATURE A ¼ preexponential factor As ¼ maximum mass loss attainable cp ¼ specific heat
CPM ¼ C/min d ¼ characteristic length scale D0.1 ¼ dispersion parameter E ¼ activation energy Eavg. ¼ average activation energy k ¼ thermal conductivity n ¼ order of reaction R ¼ universal gas constant s_ ¼ recession rate T ¼ temperature t ¼ time w ¼ instantaneous weight Xs ¼ solid conversion b ¼ heating rate d ¼ thickness of pyrolysis zone

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RUBBER–CLAY NANOCOMPOSITES BASED ON THERMOPLASTIC ELASTOMERS

a ¼ thermal diffusivity r ¼ density t ¼ characteristic time scale Subscripts DIFF ¼ f¼ MAX ¼ 0¼

diffusion final maximum initial

ACKNOWLEDGMENTS The authors would like to thank Dr. Charles Y.-C. Lee of Air Force Office Scientific Research (AFOSR) for supporting this project through Grant No. FA9550-06-10356. The authors also thank Dr. Gerry Wissler of 21st Century Polymers for his assistance in compounding and molding the polymer nanocomposite samples, Dr. Rusty Blanski and colleagues at Air Force Research Laboratory/Edwards AFB, CA for conducting the thermophysical properties measurements. We also thank Dr. Zhiping Luo at Texas A&M University/Microscopy and Imaging Center for conducting the TEM analyses.

REFERENCES 1. Woods, G. The ICI Polyurethanes Book, Wiley, New York, 1990. 2. Fabris, H. J. Advances in Urethane Science and Technology, Technomic, New York, 1976. 3. Koo, J. H. Polymer Nanocomposites: Processing, Characterization, and Applications, McGraw-Hill, New York, 2006. 4. Koo, J. H.; Pilato, L. A.; Wissler, G. E. Polymer nanostructured materials for propulsion systems. J. Spacecraft Rockets, 44(6), 1250–1262 (2007). 5. Tortora, M.; Gorrasi, G.; Vittoria, V.; Galli, G.; Ritrovati, S.; Chiellini, E. “Structural characterization and transport properties of organically modified montmorillonite/polyurethane nanocomposites”. Polymer, 43, 6147–6157 (2002). 6. Tien, Y. I.; Wei, K. H. “High-tensile-property layered silicates/polyurethane nanocomposites by using reactive silicates and pseudo chain extenders”. Macromolecules, 34, 9045–9052 (2001). 7. Tien, Y. I.; Wei, K. H. “The effect of nano-sized silicate layers from montmorillonite on glass transition, dynamic mechanical, and thermal degradation properties of segmented polyurethane”. J. Appl. Polym. Sci., 86, 1741–1748 (2002). 8. Berta, M.; Lindsay, C.; Pans, G.; Camino, G. “Effect of chemical structure on combustion and thermal behaviour of polyurethane elastomer layered silicate nanocomposites”. Polym. Degrad. Stab., 91, 1179–1191 (2006). 9. Marchant, D.; Koo, J.H.; Blanski, R. L.; Weber, E. H.; Ruth, P. N.; Lee, A.; Schaefer, C. E. Flammability and thermophysical characterization of thermoplastic elastomer

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28. UL 94: Tests for flammability of plastic materials for parts in devices and appliances. Underwriters Laboratories Inc. (UL), Northbrook, IL. 29. ASTM Standard E1354-10: Specification for heat and visible smoke release rates for materials and products using an oxygen consumption calorimeter. ASTM International, West Conshohocken, PA (2010). doi: 10.1520/E1354-10, www.astm.org. 30. Nyden, M. R.; Forney, G. P.; Brown, J. E. “Molecular modeling of polymer flammability: Application to the design of flame-resistant polyethylene”. Macromolecules, 25(6), 1658–1666 (1992). 31. Stoliarova, S. I.; Westmoreland, P. R.; Nyden, M. R.; Forney, G. P. “A reactive molecular dynamics model of thermal decomposition in polymers: I. Poly (methyl methacrylate)”. Polymer, 44, 883–894 (2003). 32. Bruns, M. C.; Koo, J. H.; Ezekoye, O. A. “Population-based models of thermoplastic degradation: using optimization to determine model parameters”. Polym. Degrad. Stab., 94, 1013–1022 (2009). 33. Galina, H.; Lechowicz, J. B. “Monte Carlo modeling of degradation of polymer networks: 2. Highly branched molecules”. Polym. Gels Netw., 6(3–4), 247–255 (1998). 34. Galina, H; Lechowicz, J. B. “Monte-Carlo modeling of degradation of polymer networks: 3. Lattice networks”. Polymer, 41(2), 615–619 (2000). 35. Galina, H; Lechowicz, J. B. “Monte-Carlo modeling of degradation of polymer networks”. Polym. Gels Netw., 6(2), 103–111 (1998). 36. Kang, S.-K.; Ku, D.-C.; Lim, J.-H.; Yang, Y.-K.; Kwak, N.-S.; Hwang, T.-S. “Characterization for pyrolysis of thermoplastic polyurethane by thermal analyses”. Macromol. Res., 13(3), 212–217 (2005). 37. Rosu, D.; Tudorachi, N.; Rosu, L. “Investigations on the thermal stability of a MDI based polyurethane elastomer”. J. Anal. Appl. Pyrol. (2010). doi: 10.1016/j.jaap.2010.07.004. 38. Rein, G.; Lautenberger, C.; Fernandez-Pello, A. C.; Torero, J. L.; Urban, D. L. “Application of genetic algorithms and thermogravimetry to determine the kinetics of polyurethane foam in smoldering combustion”. Combustion Flame, 146(1–2), 95–108 (2006). 39. Burnham, A. K.; Weese, R. K. A model-fitting approach to characterizing polymer decomposition kinetics. UCRL-CONF-203168, Lawrence Livermore National Laboratory. 40. Friedman, H. L. “Kinetics of thermal degradation of char-forming plastics from thermogravimetry application to a phenolic plastic”. J. Polym. Sci. Part C, 6, 183–195 (1965). 41. Henderson, J. B. An Analytical and Experimental Study of the Pyrolysis of Composite Ablative Materials, Ph.D. dissertation, Oklahoma State University, Stillwater, OK, 1980. 42. Ceamano, J.; Mastral, J. F.; Millera, A.; Aldea, M. E. “Kinetics of pyrolysis of high density polyethylene: comparison of isothermal and dynamic experiments”. J. Anal. Appl. Pyrol., 65, 93–110 (2002). 43. Day, M.; Budgell, D. R. “Kinetics of the thermal degradation of poly(phenylene sulfide)”. Thermochim. Acta, 203, 465–474 (1992). 44. Lee, J. C.Characterization of Ablative Properties of Thermoplastic Polyurethane Elastomer Nanocomposites, Ph.D. dissertation, The University of Texas at Austin, Department of Mechanical Engineering, Austin, TX, 2010. 45. Babrauskas, V., Grayson, S. Heat Release in Fires, Interscience Communications Ltd, London, 2009.

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46. Ho, W. K. Determination of Kinetic Parameters and Char Microstructural Analysis of Thermoplastic Polyurethane Elastomer Nanocomposites for Propulsion Systems, M.S. thesis, The University of Texas at Austin, Department of Mechanical Engineering, Austin, TX, 2007. 47. Ho, W. K.; Koo, J. H.; Ezekoye, O. A. “Kinetics and thermophysical properties of polymer nanocomposites for solid rocket motor insulation”. J. Spacecraft Rocket, 46(3), 526–544 (2009). 48. Lee, J. C.; Koo, J. H.; Ezekoye, O. A.; Lam, C. K.; Erickson, K. L. Heating rate and nanoparticle loading effects on thermoplastic polyurethane elastomer nanocomposite kinetics, AIAA Paper 2009-4096, AIAA, Reston, VA, 2009. 49. Blanksi, R.; Koo, J. H.; Ruth, P.; Nguyen, H.; Pittman, C.; Phillips, S. Polymer nanostructured materials for solid rocket motor insulation—ablation performance, Proceedings of 52nd JANNAF Propulsion Meeting, CPIAC, Columbia, MD, 2004. 50. Koo, J. H.; Marchant, D.; Ruth, P.; Nguyen, H.; Pittman, C.; Phillips, S. “Processing and Characterization of Nanostructured Materials for Solid Rocket Motors”, Proceedings of the National Space & Missile Materials Symposium, Seattle, WA, 2004. 51. Koo, J. H.; Ezekoye, O. A.; Bruns, M.; Lee, J. C. Characterization of polymer nanocomposites for solid rocket motor—recent progress, Proceedings of International SAMPE 2009 ISSE, SAMPE, Covina, CA, 2009. 52. Ho, W. K.; Koo, J. H.; Ezekoye, O. A. “Thermoplastic polyurethane elastomer nanocomposites: morphology, thermophysical, and flammability properties”. J. Nanomater. (2010), Article ID 583224. 53. Lee, J. C.; Koo, J. H.; Lam, C. K.; Ezekoye, O. A. Flammability studies of thermoplastic polyurethane elastomer nanocomposites, AIAA Paper 2009-2544, AIAA, Reston, VA, 2009. 54. Lee, J. C.; Koo, J. H.; Ezekoye, O. A. Thermoplastic polyurethane elastomer nanocomposites: density, hardness, and flammability properties correlations, AIAA Paper 2009-5273, AIAA, Reston, VA, 2009.

SECTION IV

APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES

CHAPTER 17

AUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES SAMAR BANDYOPADHYAY SUGATA CHAKRABORTY RABINDRA MUKHOPADHYAY

17.1 INTRODUCTION In the highly competitive scenario, automobile manufacturers are faced with tomorrow’s challenge to deliver reduced emissions, better safety devices’ in higher under hood temperatures, better low temperature sealing, longer warranties, extended service intervals, and stylish-comfortable interiors. They are forced to deliver all these within the best price:performance ratio. The versatility of the rubber makes it one of the most desirable materials for the automotive manufacturers. The major use of rubber in automobiles is the tire, tube, and flap. Due to rapid expansion in the synthetic elastomers, the percentage use of the elastomers in the automotive has increased dramatically. Estimation indicates that in a Mercedes E-Class car, the total amount of rubber used (excluding tires) is around 29 kg (around 1.6%, taking kerb weight of 1735 kg). In this, around 4.5 kg is general-purpose rubber (NR, BR, SBR) and rest is speciality rubber. The percentage of share is demonstrated in Figure 17.1 (sources: Indian Rubber Institute’s lecture note). Not only in car, elastomers are being used extensively in each kind of automotives starting from giant earthmovers to small bicycles. Unfilled rubber compounds are useless to prepare any automotive component. Rubber compounds are generally reinforced with fillers to achieve optimize cost performance ratio. Even today, carbon black continues to be the most important reinforcing filler in the rubber industry. But, due to its polluting nature, the ubiquitous black color of the compounded rubber material and its dependence on petroleum feedstock (for synthesis) caused researchers to look out for other “white” reinforcing agents. Clay, which has been used as cheap filler in the rubber industry, has poor Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.
2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

525

526

AUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES

2.4 kg 5.3 kg 2.0 kg 0.1 kg 1.4 kg 4.1 kg

0.5 kg 0.9 kg

3.1 kg

7.7 kg

Others FPM/FKM HNBR MVQ NBR EAM ECO ACM EPDM CR BR SBR NR

0.3 kg 0.7 kg 0.4 kg

FIGURE 17.1 The percentage of share of different types of elastomers (excluding tires) used in Mercedes E-Class car.

reinforcing ability due to its big particle size and low surface activity. For a filler to behave as a good reinforcing agent, the three main factors are particle size, structure, and surface characteristics. It was only recently that researchers succeeded in intercalating polymers into the clay layers and thereby prepared polymer–clay nanocomposites (PCNs), which not only exhibit outstanding mechanical properties but also very good barrier and thermal properties. The following sections describe the use of rubber in automotive industry as well as the potentiality of the rubber–clay nanocomposites in automotive industry.

17.2 AUTOMOTIVE APPLICATION OF RUBBER Before proceeding to the present scenario of rubber–clay nanocomposites in the automotive industry, it will be better to understand a little bit about the different elastomeric components used in the automobile. The following section deals in short with different types of rubber parts used in automobiles. Apart from tires, elastomeric products are found in many other parts of the automobile, viz. fuel systems, under hood, power train, chassis and underbody, interior trim, exterior trim, and so on. A pictorial presentation of the different rubber parts used in car is represented in Figure 17.2. Some of the major rubber products used in automotive applications are discussed below in short. The heart of any automobile, that is, engine also contains different kinds of rubber parts, which are crucial from the application point of view. Figure 17.3 shows one example of elastomers for gasoline engines reported by Akema and Yoshida [1].

AUTOMOTIVE APPLICATION OF RUBBER

527

FIGURE 17.2 Pictorial presentation to show different rubber parts in a car.

17.2.1 Automotive Hose 17.2.1.1 Brake Vacuum Hose Brake vacuum hose is designed to transmit manifold vacuum in the automotive power assisted braking system. It is generally manufactured by neoprene and reinforced with rayon. 17.2.1.2 Fuel Hose Fuel hose is designed to carry automotive fuel and is designed to operate in the temperature ranging from
30 to 100 C. It is generally made with fluoro elastomers or NBR or chlorosulfonated polyethylene (CSM). 17.2.1.3 Fuel Injection Hose Fuel injection hose is designed to carry automotive fuel and is designed to run higher pressure and temperature. The inner cover is made with PVC/NBR and outer cover is made with CSM and reinforced with fabric.

FIGURE 17.3 Elastomeric parts of a gasoline engine.

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17.2.1.4 Heater Hose Heater hose is specifically developed for vehicle cooling system It is suitable for use in the temperature range of
40 to þ100 C. It is manufactured with EPDM rubber and is twin braided spiral wrap reinforced. It is weather resistant and impervious to coolant system additives. 17.2.1.5 LPG Vapor Hose LPG vapor is specifically developed to carry LPG vapors from automotive LPG conversions. The hose is made from CR-based tube and is encased with braided cotton reinforcement. 17.2.1.6 Power Steering Return Hose This hose is specifically designed as low-pressure side return hose for vehicle power steering system. It is manufactured with NBR rubber inner reinforced with polyester and covered with a CR outer resistant of oil, fuel, and heat. Its operating temperature range is
40 to 120 C and working pressure is around 100 psi. 17.2.1.7 Break Hose Break hose is used to carry break fluid. The hose is made with EPDM (inner tube) and cover is made with CR or EPDM. The reinforcement is of braided polyester or PVA (polyvinyl alcohol) yarn (Figure 17.4). 17.2.2 Automotive Seals Seals are mechanical devices used to prevent leakage of liquid, solids, or gases. They are also used to prevent penetration of foreign particles into concealed containers or

FIGURE 17.4 Different kinds of automotive hose, seal, O-ring, and belts.

AUTOMOTIVE APPLICATION OF RUBBER

529

piping systems. Seals are available in enormous variety of design. Selection of rubber for seal depends on working pressure, temperature, corrosive atmosphere, material, shaft speed, and so on. Automotive seal can be classified as static seals (“O” rings, gasket) and dynamic seals (radial lip seals). The main requirement for the automotive seals are low compression set property, low volume swelling, low hardening or softening, and better aging resistant. Nitrile, silicon, and fluoro elastomers are extensively used for seal materials. 17.2.3 Automotive Belts The main objective of automotive belts is to transmit power from prime mover to machine or from one shaft to another. Automotive belts can be broadly divided into timing belts and V-belts or fan belts. A timing belt, timing chain, or cam belt is a part of an internal combustion engine that controls the timing of the engine’s valves. Some engines, like the flat-4 engine used in the VW Beetle, use timing gears. The term “timing belt” is also used for the more general case of any flat belt with integral teeth. Such belts are used for power transmission or to interchange rotary motion and linear motion, where either high loads or maintaining a specific drive ratio are important. A common nonautomotive application is in linear positioning systems or bicycle drives. Mainly NBR, HNBR, and CR elastomer reinforced with fabric is used to prepare the automotive belts. 17.2.4 Automotive Tubing Vacuum and washer tubing is designed to use in windscreen washer system and vacuum-operated controls. Radiator overflow tubing is suitable for use in radiator overflow application. Mainly EPDM is used for this application. It is heat and weather resistant. 17.2.5 Door Seal and Window Channels Door seals and window channels are made from EPDM rubber for better weather resistance (Figure 17.5). 17.2.6 Diaphragms and Rubber Boots Rubber boots and different types of diaphragms are prepared from EPDM and CR rubber having high ozone, chemical resistance, and excellent physical property. 17.2.7 Tire, Tube and Flap The chunk of use of elastomer is in the tire, tube, and flap. A tire is a ring-shaped covering that fits around a wheel rim to protect it and enable better vehicle performance by providing a flexible cushion that absorbs shock while keeping the wheel in close contact with the ground. The fundamental materials of

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AUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES

Designed to prevent penetration behind seal

Outside seal

Body

Door

Inside seal

FIGURE 17.5 Door seal and rubber boots.

modern tires are rubber and fabric along with other compounding chemicals. They consist of a tread and a body. The tread provides traction while the body ensures support. Before rubber was invented, the first versions of tires were simply bands of metal that fitted around wooden wheels in order to prevent wear and tear. Today, the vast majority of tires are pneumatic, comprising a doughnut-shaped body of cords and wires encased in rubber and generally filled with compressed air to form an inflatable cushion. The tires may be with tube or tubeless. A tube is a flexible, airtight hollow ring, usually made of rubber, that is, inserted into the casing of a pneumatic tire for holding compressed air. A flap is a rubber protector used in tube-type truck tires to prevent injury to the tube by the bead toes and at the valve slot of the rim. Pneumatic tires are used on many types of vehicles, such as bicycles, motorcycles, cars, trucks, earthmovers, and aircraft. The picture of pneumatic tire, tube, and flap is given in Figure 17.6.

FIGURE 17.6 Picture of tire (a), tube (b), and flap (c).

ELASTOMERIC NANOCOMPOSITES AND RUBBER INDUSTRY

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17.2.8 Other Miscellaneous Rubber Parts Besides the above-said applications, there are numerous rubber products used in automotive industry, viz. rubber bumpers, rubber engine mountings, rubber mats, grommets, rubber bellows, flexible couplings, mud flap, and so on.

17.3 PRIME REQUIREMENT OF DIFFERENT ELASTOMERIC AUTO COMPONENTS FROM APPLICATION POINT OF VIEW The general requirements of some important auto components are given in Table 17.1 [2]. The requirements are not fixed but varies based on the optimum product performance and cost. The table clearly indicates the stringent requirement of different rubber components used in the automotive industry.

17.4 ELASTOMERIC NANOCOMPOSITES AND RUBBER INDUSTRY The words “nanocomposites,” “nanomaterials,” and “nanofillers” are fairly recent but such materials were in use from the beginning of this century (e.g., carbon black is being used as a reinforcing filler in rubbers since 1904) and apparently always existed in nature (in minerals and vegetation) [3]. The major advantages that nanocomposites have over conventional composites are (i) lighter weight due to low filler loading and (ii) improved properties (includes mechanical, thermal, optical, electrical, barrier, etc.) compared to conventional composites at very low loading of filler. Three types of nanocomposites can be distinguished depending upon the number of dimensions of the dispersed particles in the nanometer range [4]. (i) Isodimensional nanofillers result when the three dimensions are in the order of nanometers, such as spherical silica nanoparticles obtained by in situ sol–gel methods [5,6] or by polymerization promoted directly from their surface [7]. (ii) When two dimensions are in the nanometer scale while the third is larger, an elongated structure result, as for example, carbon nanotubes [8,9] or cellulose whiskers [10,11], which are extensively, studied as reinforcing nanofillers yielding materials with exceptional properties. (iii) The third type of nanocomposites is characterized by only one dimension in the nanometer range. Here the filler is in the form of sheets of one to a few nanometer thick to hundreds to thousands nanometers long. Clays, clay minerals, and layered silicates belong to this family and the composites are known as polymer–clay nanocomposites (PCNs) or polymer-layered silicate nanocomposites (PLSNs).

532

Seals

Hoses

Weather strip Glass run Filler cap O-ring

Body

A/T oil cooler hose Power steering hose Brake hose Clutch hose Master vacuum hose

Chassis

Crank shaft-rear Crank shaft-front Diaphragms Valve stem seal Valve cover

Fuel hose Filler hose Air conditioning hose

Body

Engine

Fuel hose Emission control hose Air duct Water hoses Radiator, etc.

Name

Engine

Parts

EPDM EPDM NBR þ PVC

VMQ ACM ECO, NBR, FVMQ FKM NBR

ACM NBR/CR SBR/NR/CR þ EPDM SBR/NR/CR þ EPDM NBR/CR

NBR/CR NBR þ PVC NBR/CR

FKM/ECO/GECO NBR/CR, ECO, ACM NBR þ PVC EPDM Polyester

Current Materials

FKM FKM

ACM

Long-life Long-life

Heat resist.

ACM EPDM EPDM ECO

PA6/CIIR

Permeability Heat resist. Water permeability Water permeability Heat resist.

FKM/ECO/ECO

TPO, ACM EPDM (aramid)

New Materials

Permeability

Heat resist. Heat resist. Heat resist.

Needs

TABLE 17.1 General Requirements of Some of the Important Auto Components

Long-life Low abrasion

Heat resist. Heat resist.

Low permeability

Heat resist.

Low cost

Trend

533

Wheel

Others

Tube

Flap

Flap

Bumpers

Bumpers

Tube

Engine mount Mat

Chassis Mat

Tire

In-tank pump cushion Muser hanger rubber Wiper blade

Body

Tires

Synchronous belt Accessory drive belt

Transmission oil seal Power steering oil seal Ball joint dust cover Constant velocity joint boot Rack and pinion boot Brake master cup Caliper piston seal

Engine

Chassis

NR, SBR

IIR, CIIR, BIIR, EPDE

NR, SBR, BR

NR, SBR, EPDM

NR (SBR, BR) NR, SBR

NBR þ PVC EPDM CR, NR

CR CR

NBR NBR CR, U CR CR SBR SBR

Heat and cut resistance

Low growth high air retention

Tubeless tire, puncture proof tube

Fuel economy, higher mileage, longer life, low noise, high grip, and so on

Specialty SBR Lighter weight and and BR low rolling resistance, high abration, chipping and chinking resist., higher flexibility

Heat resist. Low cost Low cost, longer life, and weather resist.

Use of recycled rubber

Long-life

Long-life Heat resist.

Heat resist.

Weather resist.

Low abrasion

HNBR

TPEE TPO, TPEE EPDM EPDM

Heat, ozone Heat, ozone Heat resist., life Heat resist., life Heat resist., life

ACM

Heat resist.

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AUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES

17.5 SUPERIORITY OF CLAY/CLAY MINERAL IN COMPARISON TO OTHER NANOFILLERS There are several reasons to choose clay mineral as a potential nanofiller in comparison to other nanofillers such as carbon nanotubes, nanographite, nanoCaCO3, and so on. Clay mineral is easily available in the nature and is quite cheap in comparison to the other nanofillers. Most of the other nanofillers are required for complex synthesis process. The modification of the clay minerals or clays is easier and can be easily tailor made. Nanoclay composites demonstrate some unique properties with respect to other nanofillers. Excellent barrier property and air retention property can be obtained from clay nanocomposites. Very high elongation and antiaging property can also be obtained.

17.6 ORGANO-MODIFIED CLAY/CLAY MINERALS Montmorillonite, and other layered silicate clays, are naturally hydrophilic. This makes them poorly suited to mixing and interacting with most polymer matrices, which are mostly hydrophobic. Moreover, electrostatic forces hold the stacks of clay platelets tightly. For these reasons, the clay or clay minerals must be treated before it can be used to make a nanocomposite. These stacks of clay platelets are much larger than 1 nm in every dimension. Making a composite out of untreated clay would not be a very effective use of material, because most of the clay would be stuck inside, unable to interact with the matrix. A popular and relatively easy method of modifying the clay/clay minerals surface, making it more compatible with an organic matrix, is ion exchanging. The cations are not strongly bound to the clay surface; so small molecular cations can replace the cations present in the clay. By exchanging it with various organic cations, montmorillonite clay minerals can be compatibilized with a wide variety of matrix polymers. At the same time, this process helps to separate the clay platelets so that they can be more easily intercalated and exfoliated.

17.7 SCOPE OF APPLICATION OF ELASTOMERIC NANOCOMPOSITES IN AUTOMOTIVE INDUSTRY Today, polymer composites are widely used in automotive industries. However, these composites are fabricated by adding large amounts of microparticles, thermal stabilizers, chemical resistance, and flame resistance additives into the polymer matrix. Therefore, their improved performance often comes with the increase in materials density and low fuel efficiency. In contrast, polymer nanocomposites offer higher performance with significant weight reduction and affordable materials for transport industries such as automotive and aerospace.

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Elastomeric organoclay nanocomposites have huge scope of application in automotive industry. The following section describes the advantages of using the elastomeric clay nanocomposites in terms of mechanical, physical, thermal, air retention, aging, and processability. 17.7.1 Lighter Weight and Balanced Mechanical Property One of the primary benefits of clay/clay minerals elastomeric nanocomposites is the low specific gravity when compared to traditional fillers. This is important to note one of the key drivers in the automotive market is weight reduction. All of us ask, will greater than $3.00/gallon gasoline drive consumers to demand more fuelefficient vehicles? Couple this with the fact that consumers are demanding vehicles with more overall content, which actually adds weight, and we realize that this is working against the overall fuel efficiency of today’s vehicles. Looking forward, the opportunity for weight reduction favors materials that are lower in weight, but still provide the desired physical properties required for the specific applications. Kojima et al. [12] observed that only 10 phr (parts per hundred parts rubber) of the organoclay was necessary to achieve tensile strength comparable to an NBR compound loaded with 40 phr of carbon black. Magaraphan et al. reported that a small loading of 7 wt.% modified clay was good enough to achieve mechanical properties exhibited by high structure silica filled and carbon black filled NR vulcanizates. [13]. Joly et al. observed that organically modified interlayer of Mt (montmorillonite) were easily penetrated by natural rubber (NR) chains and led to intercalated structures along with partial exfoliation [14]. Modulus increase was comparable to that achieved by high loadings of conventional micrometer-sized fillers. Only 10 wt.% modified Mt loading demonstrated the advantages of a highsurface-area filler. Varghese and Karger-Kocsis prepared NR-based nanocomposites with 10 wt.% natural (sodium bentonite) and synthetic (sodium fluorohectorite) layered silicates by the latex compounding method [15]. Commercial clay minerals (inert material) were used as the reference material in their work. It was observed that the layered silicates outperformed the commercial clay in mechanical, thermal, and swelling tests. CR–Mt nanocomposites exhibited high hardness, high modulus, and high tear strength. Potential application areas suggested by Wang et al. for these nanocomposites were as inner tubes, inner liners, and dumpers [16]. At equivalent filler loading (20 phr), the rectorite–SBR nanocomposites exhibited better mechanical properties than SBR filled with N330 or calcium carbonate [17]. Wu et al. [18] prepared the NBR–clay nanocomposites. TEM showed that the silicate layers of clay were dispersed in the NBR matrix at the nanometer level and had a planar orientation. The NBR–clay nanocomposites exhibited excellent mechanical and gas barrier properties, which was described by Nielsen’s model. Varghese et al. prepared NR, polyurethane rubber (PUR), and NR–PUR-based nanocomposites from the related lattices by adding 10 phr pristine synthetic sodium fluorohectorite (a layered silicate) [19]. It was observed that in a blend composed of polar PUR and apolar NR, the silicate layers were preferentially embedded in the polar PUR phase in well

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intercalated/exfoliated stage. The properties of the PUR–NR-based nanocomposites were similar to those containing plain PUR. This result is of great economic significance as NR latex is cheaper than PUR latex. Nano-reinforcement was best reflected in the stiffness and strength-related properties of the rubber composites. U. S. Patent 20030144401 refers to the preparation of clay–rubber nanocomposites by the latex route and such materials have been suggested for use in tire components such as tire tread, sidewall, and/or inner liner [20]. Another U.S. Patent (2005065266) reports the preparation of nanocomposites comprised of water swellable clay particles in aqueous emulsions like anionic SBR or NR containing a novel amine for aiding in intercalation and partial exfoliation of the clay particles [21]. Applications of such rubber nanocomposites are contemplated, for example, as aircraft tire tread where significant replacement of carbon black reinforcement is desired to reduce heat buildup for tire durability and reduction in tire weight for fuel economy. Arroyo et al. prepared the NR reinforced with 10 phr Na þ -Mt (unmodified clay) and octadecylamine-modified Na þ -Mt (organoclay). They compared the results with NR loaded with 10 and 40 phr carbon black [22]. The rubber compounding was carried out on an open two-roll mill at room temperature. Vulcametric curves showed that the unmodified clay slightly changed the cure characteristics of NR and this was ascribed due to poor compatibility between the unmodified clay and hydrophobic polymer. However, the optimum cure time was sharply reduced in presence of the organoclay thereby behaving as an effective accelerating agent in NR vulcanization. The accelerating action was attributed due to the ammonium groups of the organic cations. Carbon black was also observed to decrease the optimum cure time but to a smaller extent compared to the organoclay. The nanocomposite with 10 phr of organoclay showed higher torque values than the NR mix containing 40 phr carbon black thereby indicating a higher degree of cross-linking which was also confirmed by swelling measurements and differential scanning calorimetry (DSC). Further, the organoclay behaved as an effective reinforcement agent for NR and showed a stronger reinforcing effect than carbon black while retaining the elasticity of the rubber. The mechanical behavior of NR with 10 phr organoclay was comparable to the compound with 40 phr carbon black. Some of the additional mechanical properties such as rebound resilience, hardness, compression set, and abrasion loss reported by Arroyo et al. are shown in Figure 17.7. EPDM–organophilic Mt hybrid nanocomposites were successfully prepared by a simple melt compounding process by Chang et al. [23]. The hybrid nanocomposites exhibited great improvement in tensile and tear strength, modulus, and elongation at break. Liao et al. investigated BR–organically modified montmorillonite nanocomposites along with polyisoprene and styrene–butadiene rubber nanocomposites prepared by in situ anionic intercalation polymerization [24]. The results showed that a certain extent of exfoliated nanocomposites could be prepared by in situ anionic polymerization. The incorporation of organoclay obviously changed the microstructure of BR; the concentrations of the 1,2-unit, 3,4-unit, and trans-1,4-unit increased dramatically with an increasing concentration of organoclay and the concentration of the cis-1,4 structure decreased. Organoclay apparently

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FIGURE 17.7 Comparison of compression set and abrasion loss for the various mixes. Mixes: 1, NR (unfilled); 2, NR þ 10 phr unmodified clay; 3, NR þ 10 phr organoclay; 4, NR þ 10 phr carbon black; and 5, NR þ 40 phr carbon black.

strengthened the rubber matrix; for example, the tensile strength and hardness of nanocomposites increased greatly, but the permanent deformation did not change much. Effect of different acrylonitrile content on the mechanical, dynamic mechanical, and rheological properties of nanocomposites was explained by Sadhu and Bhowmick [25–27]. The polymers with highest acrylonitrile content gave larger enhancement of properties. Acharya et al. also established structure–property relationship for EPDM–LDH nanocomposites [28,29]. They observed improvement in tensile and elongation at break property. The well-dispersed rigid modified-LDH layers efficiently transferred stress from polymer and directly enhance the stiffness in the corresponding nanocomposites. Interestingly, the elongation at break (Eb) also increased with the modified-LDH content. This increase may be due to the platelet orientation or chain slippage or plasticization. Ray and Bhowmick studied the effect of montmorillonite clay–polyacrylate hybrid material on the properties of poly(ethylene-octene) copolymer [30]. Preparation and properties of organically modified nanoclay and its nanocomposites with poly(ethylene-octene) copolymer were reported by Maiti et al. [31]. Excellent improvement in mechanical properties and storage modulus was noticed by them. The results were explained with the help of morphology, dispersion of the nanofillers, and its interaction with the rubber. Liao and Wu prepared new nanocomposites from poly(ethylene-octene) elastomer, montmorillonite, and biodegradable starch by means of a melt blending method [32]. They showed that the nanocomposites could provide a stable tensile

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strength when the starch content was up to 40 wt.%. Chang et al. produced microcellular foam from this elastomer [33]. CR–clay nanocomposites were prepared by co-coagulating the rubber latex and clay aqueous suspension and the properties of the nanocomposites were compared with conventional carbon black filled systems [16]. Cloisite 10A hybrids and the terpolymeric nanocomposites demonstrated superior mechanical and dynamic mechanical properties. Terpolymer–clay hybrids with 9 wt.% Cloisite 10A also showed higher thermal stability. Nanocomposites synthesized in situ showed better properties than those prepared in solution method. Organically modified montmorillonite nanocomposites were prepared by using dichloromethane as a cosolvent for both polymer and clay at room temperature, through the solvent-casting method. The nanocomposites showed higher tensile modulus than the polymer matrix [34]. Maiti and Bhowmick also investigated the effect of synthetic montmorillonite on the properties of fluoroelastomers [35]. The natural montmorillonite filled sample showed 65 and 51% improvement in tensile strength and 100% modulus, respectively over the control. All the synthetic clay filled samples provided better tensile strength compared to natural clay filled one. Better swelling and thermal resistance were also demonstrated by the synthetic clay based nanocomposites. Synthetic clay based nanocomposites were observed to be thermodynamically more favorable than the natural clay filled one. Star-shaped and linear block thermoplastic poly(styrene-b-butadiene) copolymer (SBS)/organophilic montmorillonite (OMt) clays were prepared by a solution approach by Liao et al. [36] The mechanical strength of nanocomposites with the star-shaped SBS/Ot were significantly increased. U.S. Patent 6,060,549 describes a method to prepare rubber toughened, clay filled nanocomposite material that can be used in automotive bumper application [37]. According to this invention, a toughened material was prepared based on a blend of a thermoplastic engineering resin (e.g., nylon), a functionalized copolymer of C4–C7 isomonoolefin (e.g., isobutylene-co-paramethyl styrene) and layered clay. The material showed superior mechanical property and enhanced impact strength. The elastomeric component was in the range of 10–30% of the total polymer content. 17.7.2 Barrier Property or Air Retention Property The pneumatic tires are typically produced so that the barrier contains the inflation air. Such barrier can be inner liner or tube or the entire tire carcass. The barrier performs critical safety and utility functions in the tire. If diffusion of the air through the tire is minimized, inflation pressure is maintained over a long period of time. Under-inflation leads to tire damage and possible catastrophic tire failure. Moreover, internal or intercarcass pressure within the tire contributes to oxidative degradation of the rubber and reinforcing fibers and to internal flaw growth during operation. One of the unique properties demonstrated by the organoclay elastomer nanocomposites is the air retention property or the barrier property. Thus, the organoclay elastomer nanocomposites has huge scope in the automotive tire inner liner, automotive tube, and so on. Better air retention property is attributed to the

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sheet-like structure with high aspect ratio of the clay. When distributed properly, the sheet-like structure of the clay create hindrance for the passage of the air. The gas barrier property of rectorite–SBR nanocomposites was excellent and higher than that of N330–SBR composite due to the higher aspect ratio of dispersed unit. Butyl rubber (IIR) and halogenated butyl rubbers are mainly used in inner tube and liner applications of tires. But both are difficult to process and hard to co-crosslink with NR, SBR, and BR, which are used in the tire body sections, due to the difference in unsaturation. Rectorite–SBR nanocomposites due to their excellent mechanical and gas barrier properties are expected to be good candidates for tire tube and inner liner applications [17]. Isobutylene isoprene rubber (IIR)–clay nanocomposites were prepared successfully by several researchers [38–42]. Kato et al. prepared them by melt processing with maleic anhydride-grafted IIR (Ma-gIIR) and organophilic clay. With the addition of 15 phr clay, gas barrier properties of the nanocomposites were 2.5 times greater than those of Ma-g-IIR [38]. A special kind of butyl rubber is brominated polyisobutylene-co-paramethylstyrene (BIMS). BIMS-based nanocomposites were prepared by Maiti et al. [43]. The nanocomposite were prepared in solution processing method using various organoclays and their mechanical, dynamic mechanical, and rheological properties were measured and explained with reference to the XRD and TEM results. The increment in barrier properties in the case of the modified clay filled BIMS was remarkable when compared with that of the gum vulcanizate. There are patents and papers on low permeability and improved barrier properties of butyl rubber nanocomposites [44–47]. Meneghetti et al. reported the SBR– organoclay nanocomposites with enhanced gas barrier property [48]. A U.S. patent [49] describes the preparation of the tire inner liner and tube compounds based on the general purpose rubber and organoclay without affecting the physical property. The invention describes the use of a reactive rubber to prepare organoclay nanocomposites, which can be used in tire inner liner. The reactive rubber must have a chemical group capable to form onium ions, which in turn will react with the dispersed clay to form nanocomposites. Another patent [50] describes the use of BIMS and layered silicate to prepare nanocomposites. A method has been described to prepare a clay–polymer nanocomposites via an organic aqueous emulsion wherein the polymer is provided in the organic phase, and the clay is provided in the aqueous phase. Noticeable improvement in the permeation rate has been observed. The nanocomposites can be used in tire inner liner and tube compounds. There are numbers of patent describing the use of organoclay elastomer nanocomposites in tire inner liner and tube [51–53]. 17.7.3 Aging and Ozone Resistance Elastomeric composites are very much prone to environmental ozone and oxygen attack. They are easily attacked by oxygen and ozone due to the presence of unsaturation in the backbone. To protect the rubber components, antioxidant and antiozonate are added in the formulation. However, authors have demonstrated that

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in situ resol modified bentonite clay nanocomposites exhibited better aging property in comparison to regular compound [54]. Authors have investigated and found that organoclay filled SBR nanocomposite exhibited better retention of physical property after thermal or thermooxidative aging. Better performance of organoclay filled compound was probably due to slow release of the adsorbed curatives as well as the barrier property of the compound. The exfoliated and intercalated clay layers slowed down the diffusion of the oxygen inside the rubber matrix. Thus, the degradation process was less severe in organoclay filled compound. It was also observed that ozone resistance of the organoclay filled compounds were better in comparison to carbon black filled compound. The superior barrier property of the dispersed organoclay (Cloisite 20A) in compound (OC), create hindrance for the ozone attack on the surface and subsequent stages. Thus, compound OC exhibited better ozone resistance property in comparison to carbon black filled compound (CB). The ozone cracked sample picture is shown in Figure 17.8. Figure 17.9 represents the relative physical properties of the compounds after thermo and thermooxidative aging. For determination of accelerated thermooxidative aging property, tensile and hardness specimens were air aged at 105 C for 7 days in a multicell aging oven. For accelerated thermal aging, the samples were aged inside the compression mould. The molding time was increased to 30 h for 130 C.

FIGURE 17.8 Ozone crack photograph of compound OC (A1, 24 h and A2, 48 h) and compound CB (A1, 24 h and A2, 48 h) at 30 magnifications.

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FIGURE 17.9 Relative decease/increase of physical properties after accelerated thermal and thermooxidative aging. (“ þ ” sign indicates relative increase, “
” sign indicates relative decrease, “Mod” stands for 300% modulus, “Ts” stands for tensile strength, “Eb” stands for elongation at break, “Hd” stands for hardness, “T” stands for thermal aging, and “TO” stands for thermooxidative aging.)

The aging performance of HNBR–clay nanocomposites was studied at 178 C in air, water, and oil, respectively [55]. It was clear that adding organoclay significantly improved material performance at high temperature in both air and oil. HNBR nanocomposites maintained about 80% of tensile strength after aging in air and 50% of tensile strength after aging in oil. However, HNBR reinforced with conventional carbon black lost most of its mechanical properties after aging in air and oil. Aging performance of HNBR nanocomposites in water was similar to that of carbon black filled HNBR. This was may be due to the swelling of the material with the presence of organoclay, which counter-interacts with its reinforcing effect. It was also found that the retention of material properties was not affected much by clay content when it is greater than 2.5 phr. The addition of organoclay in the rubber compound can improve the thermooxidative aging property and the ozone resistance property besides other physical properties. Thus, the organoclay containing nanocomposites can be used in different automotive applications such as door–window sealing, automotive bumpers, and others. 17.7.4 Solvent Resistance Most of the under hood automotive runner product requires the solvent resistance (rubber hose, tube, etc.) besides other physical property. Elastomer clay nanocomposite exhibits improvement in the solvent resistance property. Maiti and Bhowmick investigated the effect of synthetic montmorillonite on the properties of fluoroelastomers [56]. Better swelling and thermal resistance were also demonstrated by the synthetic clay based nanocomposites. Synthetic clay based nanocomposites were observed to be thermodynamically more favorable than the natural clay filled one. It has been reported that the addition of small amount of clay reduced the swelling of the NR compound [57]. Solvent up take measurement of the thermoplastic elastomeric clay nanocomposites vulcanizate was conducted in

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toluene [58,59]. Both the room temperature and 100 C result indicated that the solvent up take decreased with clay content. Maiti and Bhowmick also investigated the diffusion and sorption of methyl ethyl ketone and tetrahydrofuran through fluoroelastomer–clay nanocomposites in the temperature range of 30–60 C by swelling experiments [59]. The overall sorption value decreased with the addition of the nanoclays and was maximum for the unmodified clay filled sample, which also demonstrated the slowest rate of increase of solvent uptake. As the temperature of swelling was increased, the solvent uptake increased in all the systems. Kim et al. further studied the morphological, rheological, and swelling properties of nanocomposites based on NBR and organophilic layered silicates using the melt processing [60]. Vulcanized neat NBR and organo-Mt–NBR composites were immersed in methyl ethyl ketone (MEK) and the extent of swelling clearly decreased with increase in organo-Mt loading. Thus, the organo-Mt–NBR nanocomposites have excellent barrier properties compared with vulcanized NBR. Addition of layered nanoclays to a neat polymer restricted the permeability of nanocomposites and reduced it from that of the neat polymer by the product of the decreased area and the increased path length. When a solvent diffuses across a neat polymer, it must travel the thickness of the sample. When the same solvent diffuses through a nanocomposite film with nanoclays, the distance it must travel around each clay layer it strikes increases its path length. Thus, the nanocomposites has very high potential to be used in the automotive hoses and tubing due to its inherent solvent resistance property. 17.7.5 Better Processability Success of any new technology depends on the adoption of the same by both the manufacturer and the end user. It is very much applicable for the automotive component manufacturers. Besides better performance, ease in processability is the key factor for the automotive component manufacturers. Bhowmick et al. [61] reported that in case of NBR–organoclay nanocomposites, shear viscosity continuously decreases with increasing shear rate, that is, shear thinning effect occurred. The extrudate profile was also better in case of nanocomposites. The investigations illustrated that the processability improved on incorporation of the organoclay. Figure 17.10 shows the SEM images of the extruded compounds. The melt rheology of ACM (poly acrylic rubber)/silica and ENR (epoxidized natural rubber)/silica hybrid nanocomposites has been studied in a capillary rheometer at 100, 110, and 120 C and at nine different shear rates [62]. It was found that the shear viscosity does not increase to a great extent for the rubber–silica nanocomposites even on substantial increase in the concentration of nanosilica. This has been a striking feature of these systems and is not encountered with ordinary filled composites usually. Both the nanocomposites show pseudoplastic behavior under the experimental temperatures and follow power law model. The “k” value for the nanocomposites is not significantly higher compared to the gum and this value

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FIGURE 17.10 SEM image of the extruded compounds (Compound (i) is the gum NBR compound (34% ACN containing) having no clay, compound (ii) contains 4 phr of sodium montmorillonite, and compound (iii) contains 4 phr of organo-modified montmorillonite clay) [61].

decreases with increasing extrusion temperature. ACM–silica nanocomposites show lowering in die swell with shear rates up to some critical shear rates, beyond which the die swell increases. ENR–silica nanocomposites exhibit increasing die swell at higher shear rates. Both the nanocomposites display decrease in die swell with the increase in the nanofiller content as well as the temperature. The extrudate roughness decreases with the generation of nanosilica that indicates higher viscous nature in the nanocomposites compared to the gum rubbers. It is also supported from the lowering of first normal stress differences in the specimen at particular temperature. The distribution of silica is almost uniform after the nanocomposites being extruded out of the capillary. As a whole, the ENR–silica nanocomposites displayed regular flow compared to ACM/silica system under the experimental conditions owing to better polymer–silica interface in the former. The activation energy of the flow shows a

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TABLE 17.2 Cure Characteristics of Different Compounds Material

A4160

A5140

A5160

A6140

A6160

A7140

A7160

Maximum torque (Tmax), (dN m) Minimum torque (Tmin), (dN m) DTorque ¼ Tmax
Tmin, (dN m) TS02 (min) TC90 (min) CRI (min–1)

6.17 0.9 5.3 11.09 22.85 8.5

7.9 1.0 6.9 12.15 24.48 8.1

6.5 0.8 5.7 3.59 10.26 15.0

7.7 1.0 6.7 19.95 33.50 7.4

6.5 0.9 5.6 5.3 12.1 14.7

7.6 1.0 6.6 24.94 40.61 6.4

6.6 0.8 5.8 6.48 12.89 15.6

Nomenclature: AXT, particular compound cured at specific temperature T; A4, ordinary clay, control compound; A5, organoclay; A6, organoclay and 0.25 phr CTP; A7, organoclay and 0.5 phr CTP. All the compounds contain 5 phr organoclay.

decreasing trend with the increase in shear rates (up to a certain shear rate) for both the hybrid nanocomposites. It increases almost linearly with the silica content at a particular shear rate. It has been reported that the organoclay nanocomposites act as curing activator. The amine moiety of the organoclay forms a complex with the accelerator/activator. The complex reduces the activation energy of the vulcanization reaction [63]. A rheometric study was carried out by the authors (unpublished work) at 140 and 160 C on SBR–organoclay nanocomposites. The rheometric property is given in Table 17.2. It was clear from Table 17.2 that 20 C reduction of curing temperature brought the cure rate index, TS02 and TC90 of compound A5 very close to the control compound (A4). However, curing at 140 C, the CRI (cure rate index or rate of curing reaction) of A5, A6, and A7 was nearly half to that of high temperature curing at 160 C. The effect of the CTP was also clear from the table. The TS02 of the A7160 is nearly double to that of the compound A5160. Thus, the addition of CTP helped to increase the scorch safety of the compounds. Manchado et al. [64] reported that the addition of organoclay reduced the activation energy of the NR–organoclay nanocomposites. Thus, it can be concluded that the reduction of the curing temperature can match the curing parameter of the amine-modified clay-containing compound to that of the control compound. Moreover, the reduction in the curing temperature is cost effective. Addition of CTP was also found to control the curing parameters. Thus, the organoclay nanocomposites is helpful to reduce the production cost either by reducing the temperature or by reducing the cycle time for production. 17.7.6 Elastomeric Polyurethane–Organoclay Nanocomposites TPUs are a member of the broader class of thermoplastic elastomer (TPE) materials. TPEs combine the flexibility and resilience of rubbers with the processability of plastics. TPU materials currently comprise approximately 15% of the volume of TPE sold annually. Therefore, it is necessary to discuss the elastomeric polyurethene–organoclay nanocomposites separately.

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Some common automotive applications for TPEs include automotive interiors, bumpers, covers, trim, constant velocity joint boots, and so on. The various TPE classes compete on the basis of flexibility, ultimate properties, adhesion, and compatibility with other part components, chemical resistance, and price. The effects of organoclays on the properties of PU were studied [65–80]. Wang and Pinnavaia prepared PU nanocomposites by solvation of organoclay by polyol first. Loading of polyol with clay up to 10–20 wt.% makes a pourable mixture [66]. XRD demonstrates that intercalation of polyol into clay results in an increase of basal spacing from 1.8–2.3 nm to 3.2–3.9 nm. Such spacing testifies of intercalation of polyol into clay. Formation of PU results in further increase of basal spacing up to more than 5 nm. The latter case may be considered as exfoliation of a clay or dispersal of nanolayers. Important, that onium ions of the clay were considered as active reagents for coupling with diisocyanate. Loading of PU with 5–10 wt.% of clay results in a two to three times improvement of tensile properties of a polymer, namely increase of strain-at-break, tensile modulus, and tensile strength. Common inorganic fillers are commonly used in PU chemistry to reduce formation cost and to increase stiffness, but the improvements in modulus for conventional PU composites are compromised by a sacrifice of elastomer properties. The nanocomposites reported in Ref. [66] exhibited an improvement in both elasticity and tensile modulus. Clay nanolayers, even when aggregated in the form of intercalated tactoids, strengthen, stiffen, and toughen the matrix in the studied case. The enhancement in strength and modulus is directly attributed to the reinforcement provided by the disperse clay nanolayers. The improvement in elasticity is tentatively attributed to the plasticizing effect of onium ions, which contribute to dangling chain formation in the matrix, as well as to conformational effects on the polymer at the clay–matrix interface. Millable PU–organoclay nanocomposites was prepared by Mishra et al. [81]. WAXD result suggests intercalation of the polymer inside the organoclay. Improvement in tensile strength, 100% modulus, and elongation at break were found. McKinley et al. [72,82] incorporated unmodified, low aspect ratio hectorite (LaponiteTM) into a number of TPUs at up to 20 wt.% using a novel solvent exchange technique. They reported remarkable mechanical property, heat distortion temperature, and morphological changes in fully delaminated nanocomposite systems. Preparation, characterization, mechanical and barrier properties, morphology and effect of processing conditions have been reported on polyurethane-based clay nanocomposites [83–93]. Igor et al. [92] showed that the coefficient of thermal expansion of polyurethane elastomers increased by 85% with the addition of 10 wt.% of nanoclay, and by 98% with 15 wt.% of carbon nanofibers. 17.7.7 Use of Organoclay Nanocomposites in Tire The major potential field of application of elastomeric nanocomposites is the automotive tire. There are lots of patents in this field. A brief discussion of some selected patents is given below to explain the usefulness of the technology. A U.S. patent [94] demands that the use of organoclay has reduced the tan d value at 1, 5, and 10% strain in comparison to carbon black filled compound. Besides, it has

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also shown higher storage modulus value, which indicates higher stiffness value of the compounds. Lower tan d value indicates lower rolling resistance, which in turn indicates better fuel economy. It has also been demonstrated that the use of in situ generated elastomer–clay nanocomposites generated by latex blending technique, further improved the result. Claim has been made for the use of this type of nanocomposites in tire tread. The experiment was carried out in RPA at 100 C. Another patent [95] describe the use of Vinyl pyridine–styrene–butadiene ter copolymer type rubber and Cloisite 30B organoclay nanocomposites. It has been claimed that the elastomer-Cloisite 30B-carbon black combination exhibited superior tan d value and the storage modulus value according to RPA measurement. Higher tensile and elongation at break have also been claimed. Noticeable improvement in the tan d value was found when organoclay was used along with coupling agent (Si69). Also, around 15 parts of N299 carbon black was replaced by organoclay and coupling agent. U.S. Patent 006,759,464B2 [96] describes the preparation and use of organoclay–elastomeric nanocomposites in tire tread, sidewall, and inner liner. In situ generated organoclay nanocomposites were used for compounding. Functionalized elastomers, such as epoxidized polyisoprene or polybutadiene rubber latex were used for the preparation of the nanocomposites master batch. Improvement in the strength and tan u value has been claimed. Elastomeric organoclay nanocomposites can be used in bead filler, sidewall insert, tread base or tread under layer [97]. The patent said that the use of epoxidized rubber along with layered clay material improved the property. Apart from the epoxidized rubber, a base polymer has also been used as matrix. The patent also described the preparation of the epoxidized rubber and the nanocomposites. Similar invention [98] claimed the use of the organo-modified montmorillonite clay nanocomposites in the tire apex or sidewall insert compounds. The invention described that the use of only 5 parts of organoclay with 55 parts of carbon black in a natural rubber based compound increased the stiffness on the compound (storage modulus) with marginal increase in tan d value. It has also been shown that the marginal drop in ultimate tensile strength and elongation at break with large gain in modulus value. Rubber toughened thermoplastic nanocomposite materials have been invented [37] and it can be used for automotive bumper application. The material not only exhibited superior tensile and modulus property but also improved impact strength. The materials were prepared based on a blend of one or more thermoplastic engineering resins, for example, nylon, a functionalized copolymer of a C4–C7 isomonoolefin; for example, isobutylene, and a para-alkylstyrene; for example, para-methylstyrene, and further contained a uniformly dispersed exfoliated layered clay-like montmorillonite. U.S. Patent 6,818,693B2 [99] described the use of the organoclay nanocomposites in racing tire tread compounds. The compound formulation is given in Table 17.3. The properties are summarized in Table 17.4. It is observed from Table 17.4 that the hardness of the carbon black filled compounds is higher in comparison to the clay-filled compounds. Moreover, the tan d values of the clay-filled compounds are higher compared to the carbon black

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TABLE 17.3 Compound Formulation [99] Materials (phr)

A

B

C

Natural rubber Polybutadiene rubber SBR 1500 Carbon black (N220) Montmorillonite clay modified with dimethyldioctadecyl ammoniums ions Aromatic oil Stearic acid Zinc oxide Antioxidant Sulfur Accelerator

60 40 – 50 –

60 40 – – 50

– – 100 80 –

5 2 3 1 2 2.5

5 2 3 1 2 2.5

D – – 100 – 80

40 2 3 1 2 2.5

40 2 3 1 2 2.5

TABLE 17.4 Physical Properties of the Compounds [99] Property Tensile strength (MPa) 300% Mod (MPa) Tear strength (MPa) Shore A hardness Rebound at room temperature (%) Rebound at 70 C (%) Tan d, 0 C Tan d, 30 C Tan d, 60 C

A

B

C

D

20.0 15.04 31.7 67.5 58.1 64.5 0.109 0.065 0.061

9.2 5.35 21.1 66.6 48.7 38.1 0.160 0.248 0.210

15.2 6.83 27.8 53.3 47.0 60.6 0.278 0.237 0.194

3.6 2.35 10.7 42.7 17.0 29.2 0.393 0.433 0.334

filled compounds. A higher tan d value indicates the higher friction coefficient and this, in turn, is a gauge for the skid resistance and or the grip of the tire on the road. The compounds according to the invention are particularly suitable for tires requiring strong grip on road such as racing tires. It has been claimed that the organoclay nanocomposites can be used for the color tire sidewall compounds [100]. The compound formulation is given in Table 17.5. Physical properties are summarized in Table 17.6. Higher strength property was observed for the nanocomposites. Better ozone resistance was also found for the nanocomposites. Mechanism for the better ozone resistance has been explained earlier (Section 17.7.3). Similarly, U.S. Patent 20030144401 refers to the preparation of clay–rubber nanocomposites by the latex route and such materials have been suggested for use in tire components such as tire tread, sidewall [20]. Another U.S. Patent (2005065266) [21] reports the preparation of nanocomposites, which can be used in the aircraft tire tread where significant replacement of

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TABLE 17.5 Compound Formulation [100] Materials (phr) Natural rubber (SMR 5) Chlorobuytl rubber EPDM rubber Silica Titanium dioxide Plasticizer Wax Nanocomposites master batch (35 phr natural rubber)

Sample A

Sample B

35 50 15 20 30.2 5 3 –

0 50 15 20 30.2 5 3 40

With same curing package.

TABLE 17.6 Physical Properties of the Compounds [100] Properties Tensile strength (MPa) Elongation at break (%) 300% Mod (MPa) Hardness, Shore A Rebound, RT (%) Dynamic ozone test (0–60% strain, 48 h, 50 pphm, 40 C) Taber abrasion test, wt. loss after 1000 cycles (g)

Sample A

Sample B

14.7 800 2.6 48 35 Broken, edge crack 1.12

11.9 607 4.8 56 35 No crack 0.86

carbon black reinforcement is desired to reduce heat buildup for tire durability and reduction in tire weight for fuel economy. Lower abrasion loss was also claimed for the compounds containing optimum loading of nanoclay ( 9 phr). U.S. Patent 200440147661 A1 [101] claimed both higher coefficient of friction and low abrasion loss of the natural rubber organoclay nanocomposites. There are numbers of sillier patents [102–106] describing the use of the organoclay elastomer nanocomposites in the tire. 17.8 DISADVANTAGES OF USE OF ORGANOCLAY ELASTOMERIC NANOCOMPOSITES IN AUTOMOTIVE INDUSTRY The main constrain for the use of organoclay elastomeric nanocomposites in automotive industry is the “cost.” Besides cost, there are some disadvantages of organoclay nanocomposites. Unlike carbon black, the nanoclay composites exhibit very high compression or tensile set property [17,107]. The high set property restricts the use of the organoclay in tire tread, tube, and different kinds of seals where the low compression set property is required. Apart from the set property, organoclay nanocomposites exhibit low fatigue life in comparison to the conventional carbon black filled compounds [108].

CONCLUSION

549

The better fatigue life of carbon black filled compounds can be explained based on the relatively better rubber to filler interaction [109]. Carbon black is nanoparticles with more or less spherical nature. The interaction of the carbon black with rubber is a strong physicochemical interaction. However, the interaction of the clay with rubber is due to van der Waals force of interaction. The interaction at the edge of the clay particles is the weakest. Therefore, during the cyclic deformation the stress concentration is highest at the edge of the clay particles. Thus, the flaws can generate easily at these points in comparison to the strongly attached nearly spherical carbon black particles. Thus, the organoclay-containing compounds exhibited low fatigue life compared to carbon black filled compound. Lower fatigue life, reduces the chances of use of the organoclay in dynamic automotive elastomeric components. However, treatment of the organoclay with silane improved these properties [109]. Same patent [108] also reported that the addition of the organoclay nanocomposites produced relatively low room temperature and high temperature self–self/ self–carcass adhesion. Low adhesion is a serious concern related with the product performance. Lower adhesion will cause the early failure of the product.

17.9 CONCLUSION Unlike plastics, probably there is not a single organoclay elastomeric nanocomposites is available commercially in the marker. Although, there are a number of advantages of using the organoclay elastomer nanocomposites, still there are few problems, which need to be addressed. The plastic parts in automotive industry is mainly for static applications whereas, the rubber parts are built for dynamic application. The lower dynamic property is the main disadvantage for nanoclay elastomer nanocomposites. The thermosetting nature of the product itself restricts the industrial trial of the organoclay elastomer nanocomposites. In thermoplastics field, even if the organoclay composites does not meet the requirements, it can be further reused in noncritical applications by remelting and reprocessing. Although, the organoclay elastomer nanocomposites has a number of advantages in terms of processability, still it requires intense (time) mixing that increases the cycle time for production. However, use of organoclay master batch can solve this problem. A lot of publications indicate that the latex stage blending, and preparation of the nanocomposites master batch could be a very promising route [110–112]. There are numbers of patent and publications in the organoclay elastomer nanocomposites. Still, there is not a single organoclay elastomeric nanocomposite available in the market for automotive application. Of course, there are few thermoplastic nanocomposites available commercially for automotive application (FORTETM nanocomposites from NOBLE POLYMERS, NYLON 6 nanocomposites from UBE INDUSTRY, NYLON 6 nanocomposites from TOYOTA MOTORS, etc.). Even though, not a single commercially available elastomer clay nanocomposites is there for automotive application, still there are number of advantages for the clay nanocomposites. Substantial improvement in physical property can be obtained with

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low level of organoclay loading compared to highly carbon black loaded compound. This will reduce the overall weight of the tire, which in turn will definitely create a positive impact on the fuel economy. Low fifer loading will also reduce the heat build up in the tire, that is, a “cooler running tire.” Very good air retention property makes the nanocomposites suitable for the inner liner and automotive tube. Inherent solvent resistance property of the clay nanocomposites makes it suitable for use in automotive hose or inner liner application. Lower abrasion loss of the clay nanocomposites makes it suitable for the tire application in term of better mileage. Better ozone and aging property makes the clay elastomer nanocomposites suitable for the tire sidewall application. A higher tan u value of the clay elastomer nanocomposites indicates higher friction coefficient and this, in turn, is a better candidate for the tires having higher skid resistance and or the grip of the tire on the road. Besides, good mechanical strength and processability makes it suitable for production of different automotive parts. In a nutshell, it can be stated that the organoclay elastomer nanocomposites is a new technology for automotive industry, which needs to be refined. It needs much more attention from the automotive parts manufacturers to short out different problems associated with it. More systematic and intense research is required to make this powerful technology suitable for automotive industry. Fortunately, big automotive giants and chemical companies are stepping into the systematic research in the field of organoclay elastomer nanocomposites. ACKNOWLEDGMENT The authors thank HASETRI and JK Tyre management for kind permission to publish this work. REFERENCES 1. Akema, T.; Yoshida, H. Elastomers in automobile. Polyfile, 25, 24–36 (1987). 2. Hashimoto, K.; Maeda, A.; Hosoya, K.; Todani, Y. Specialty elastomer for automotive applications. Rubber Chem. Technol., 71, 449–519 (1998). 3. Donnet, J.B.; “Nano and microcomposites of polymers elastomers and their reinforcement”. Comp. Sci. Technol., 63, 1085–1088 (2003). 4. Alexandre, M.; Dubois, P. “Polymer-layered silicate nanocomposites: preparation, properties and uses of a new class of materials”. Mater. Sci. Eng. Reports, 28, 1–63 (2000). 5. Mark, J. E. “Ceramic-reinforced polymers and polymer-modified ceramics”. Polym. Eng. Sci., 36, 2905–2920 (1996). 6. Wen, J.; Wilkes, G.L. “Organic/inorganic hybrid network materials by the sol–gel approach”. Chem. Mater., 8, 1667–1681 (1996). 7. Von Werne, T.; Patten, T. E. “Preparation of structurally well-defined polymer
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96. Ajbani, M.; Geiser, J. F.; Parker, D. K. U.S. Patent 006,759,464B2 (to Goodyear Tyre & Rubber Company) (2004). 97. Caprio, M.; Galimberti, M.; Grassi, L.; Testi, S.; Tirelli, D. U.S. Patent 0,065,117A1 (to Pirelli Pneumatic S.P.A) (2009). 98. Larson, B. K. U.S. Patent 006,598,645B1 (to Goodyear Tyre & Rubber Company) (2003). 99. Li, D.; Peiffer, D. G.; Elspass, C. W.; Wang, H. U.S. Patent 006,060,549 (to Exxon Chemical Patent Inc.) (2000). 100. Zhao, J.; Crawford, M. J.; Puhala, A. S.; Cohen, M. P. U.S. Patent 20060135671A1 (to Goodyear Tyre & Rubber Company) (2006). 101. Yaakub, A.; Kuen, C. P.; Keane, N.; Ross, M. U.S. Patent 20040147661A1 (2004). 102. Elspass, C. W.; Peiffer, D. G.; Kresge, E. N.; Hsieh, D.; Chludzinski, J. J.; Liang, K. S. U.S. Patent 005,883,173 (to Exxon Research and Chemical Company) (1999). 103. Weng, W.; Dias, A. J.; Karp, K. R.; Johnston, M. W.; Gong, C.; Neagu, C.; Poole, B. J. U.S. Patent 20070129477A1 (to Exxon Mobil Chemical Company) (2007). 104. Yang, X.; Cohen, M. P.; Senyek, M. L.; Parker, D. K.; Cronin, W.; Lukich, L. T.; Francik, W. P.; Gurer, C. U.S. Patent 007,342,065B2 (to Goodyear Tyre & Rubber Company) (2008). 105. Parker, D. K.; Larson, B. K.; Yang, X. U.S. Patent 20040054059A1 (to Goodyear Tyre & Rubber Company) (2004). 106. Heinrich, G.; Herrmann, W.; Kendziorra, N.; Pietag, T.; Recker, C. U.S. Patent 20020095008A1 (to Continental Aktiengesellschaft) (2002). 107. Wang, Y.; Zhang, L.; Tang, C.; Yu, D. “Preparation and characterization of rubber–clay nanocomposites”. J. Appl. Polym. Sci. 78, 1879–1883 (2000). 108. Dias, A. J.; Jones, G. E.; Tracey, D. S.; Waddell, W. H. U.S. Patent 20040132894A1 (to Exxon Mobil Chemical Company) (2004). 109. Chakarborty, S.; Sengupta, R.; Dasgupta, S.; Mukhopadhyay, R.; Bandyopadhyay, S.; Joshi, M.; Ameta, S. C. “Effect of treatment of bis(3-triethoxysilyl propyl)tetrasulfane on physical property of in situ sodium activated and organomodified bentonite clay—SBR rubber nanocomposite”. J. Appl. Polym. Sci. 116, 1660–1670 (2010). 110. Karger-Kocsis, J.; Wu, C.M. “Thermoset rubber/layered silicate nanocomposites. Status and future trends”. Polym. Eng. Sci., 44, 1083–1093 (2004). 111. Zhang, H.; Wang, Y.; Wu, Y.; Zhang, L.; Yang, J. “Preparation of reactive mineral powders used for poly(sodium acrylate) composite superabsorbents”. J. Appl. Polym. Sci. 97, 844–861 (2005). 112. Wu, Y. P.; Zhang, L. Q.; Wang, Y. Q.; Liang, Y.; Yu, D. S. “Structure of carboxylated acrylonitrile–butadiene rubber (CNBR)–clay nanocomposites by co-coagulating rubber latex and clay aqueous suspension”. J. Appl. Polym. Sci., 82, 2842–2848 (2001).

CHAPTER 18

NONAUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES CARRIE A. FEENEY HARRIS A. GOLDBERG

18.1 WATER-BASED NANOCOMPOSITES 18.1.1 Barrier Properties The concept of using nanodispersed clay minerals in polymers to create a tortuous path and thus improve their barrier properties has been known for a long time. Most of that work has used exfoliated montmorillonite. Honeywell and others have demonstrated up to a fourfold improvement in the barrier properties of nylon by the addition of nanoclays [1–3]. Although these were commercialized, their cost/performance benefit was modest and they are not in wide use. Nanotechnology—that is, the control of materials and structure on the nanometer length scale—has been central to the development of an elastomeric barrier coating technology platform [4–7], and to its high performance barrier coating products. The barrier coating technology platform starts with clay minerals that are exfoliated and dispersed in water. These materials have a plate thickness of
1 nm, and a lateral dimension of 1–30 mm in the case of vermiculite. They are used to form a tortuous path in our polymer-based coatings. The nanometer thickness is critical to our ability to have a large number of plates in a thin (
10–30 mm) coating, and to the relevance of the tortuous path models discussed below. The second important nanotechnology aspect of the barrier coating technology is the ability to control the interface between the nanodispersed clay and the polymer. Most other work on clay–polymer nanocomposites control this interface by ion exchange of the sodium ions on the surface of the clay with anionic surfactants that make the clay surface hydrophobic and improve its interaction with the polymer. Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti.  2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

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The elastomeric barrier coating technology platform does not take that approach, as it would make the clay come out of water suspension and make it impossible to have a stable water-based coating. Instead, the interface is controlled through other proprietary treatments of the clay surface, and by the addition of cationic surfactants, dispersing aids, and other additives. The goal is to create an interface that is stable, has low permeability, and provides adequate mechanical properties for the coating. In these elastomeric nanocomposites, it is the control of this interfacial interaction that enabled the coating to remain flexible even in high temperature, very dry environments. This is a critical requirement for a coating in tires and most other elastomeric products. Finally, it is critical that the nanoclays remain well dispersed in the polymer even as the coating dries. The fact that this can be achieved was surprising but has been well documented in a collaborative research effort with Professor Don Paul’s group at the University of Texas. In the sections below, each of these four components of the elastomeric barrier coating nanotechnology are addressed. 18.1.1.1 Dispersion and Orientation in Water-Based Dispersions The elastomeric nanocomposite coatings are made by combining an aqueous dispersion of butyl rubber with an aqueous dispersion of nanodispersed exfoliated vermiculite. The aspect ratio of the vermiculite is extremely high (
10,000). That is because the plate thickness is 1–3 nm, while the lateral dimensions are
10 mm (although many particles that are in the 1–10 mm range are observed). The 1–3 nm size and very large aspect ratio of the vermiculite utilized is easily demonstrated by the locally ordered domains that form even at very low concentrations in water. Figure 18.1 below shows a suspension of Microlite in water at a concentration

FIGURE 18.1 0.02 wt.% of Microlite in water showing oriented domain structure. Left: between crossed polarizers; right: with additional optical compensator.

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of
2  10  4 (0.02%) when viewed through cross polarizers. The structures are macroscopically oriented domains. Since the critical concentration for mesoscopic behavior is
1/aspect ratio, this observation confirms the high aspect ratio for the vermiculite nanoclays. It also confirms the nanodispersed nature of the exfoliated vermiculite in water. The mesoscopic orientation of the nanoclays makes it easy to achieve excellent filler orientation in thin coatings, an important feature to achieve low permeability coatings. This mesoscopic structure in the liquid phase will be even more pronounced at the much higher concentrations actually used in our commercial formulation. It is hard to observe it directly as the polymer dispersion generally interferes with getting light through the sample, although indirect evidence for it in acoustic measurements has been observed. The nanosized clay particles are important not only to achieve oriented nanocomposites but also for the stability of the water-based elastomeric formulations. This results in the shelf life of Air D-Fense 2000 at a minimum of 4 years. Although one needs to stir the formulation before use, there is no hard settling and the barrier properties remain constant even after years of storage. This stability is also due to the detailed attention paid to total solid content and the ionic concentrations as discussed in Ref. [8]. 18.1.1.2 Nanoclay Dispersion—Direct Observation of Filler Orientation The incorporation of plate-shaped nanodispersed clay particles has been the focus in order to reduce the permeability of polymer coatings. The primary function of the nanoclays is to create a tortuous path, which significantly reduces the permeation rate through the polymer. There has been significant work on modeling the effect of fillers and the tortuous path on the permeation rate. In those models, any change in the intrinsic diffusion constant and solubility of the polymer is ignored. The absolute dimensions of the particles are typically not as important as their shape. For plateshaped particles, the aspect ratio is the length (or width) divided by the thickness. Before discussing the models in detail, we will first review the direct observations of orientation and nanodispersion that have been made on the elastomeric barrier coating nanocomposites discussed previously. Electron Microscopic Observations In collaboration with Professor Don Paul and his group at the University of Texas, a significant effort was made to examine the dispersion of the vermiculite plates in the butyl matrix. The following discussion is taken directly from the resulting publication [9]. Figure 18.2 shows SEM micrographs of edge views of cryofractured butyl rubber membranes containing 20 and 30% by weight of vermiculite. Owing to the irregular nature of the fracture surface and the relatively low magnification, it is difficult to say much about the local morphology of the composites. However, it is quite apparent that the vermiculite layers are well aligned within the plane of the membrane as expected based on the method of preparation outlined earlier. Figure 18.3 shows representative TEM micrographs taken of microtomed crosssections from butyl rubber membranes containing 20 and 30% by weight of

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FIGURE 18.2 SEM micrographs of butyl rubber nanocomposites.

vermiculite. These represent but a few of many attempts to image the morphology of these composites. It is important to understand that it is very difficult to cleanly cut sections of these materials because of the very low glass transition temperature, about  80 C, of the butyl rubber matrix. The thin sections were invariably wrinkled and otherwise distorted. These difficulties stem from ensuring that the butyl rubber matrix is indeed below its Tg, and therefore rigid, at the surface being cut plus the distortions that arise from warming to room temperature which relieves stresses induced by cutting. Given these limitations, it appears that these composites consist of a mixture of individual vermiculite platelets and some stacks of platelets. The welldispersed platelets appear to be 1–2 nm in thickness and 200–400 nm in length. The thicker stacks are several micrometers in length. The wavy or curved-type structures seen in Figure 18.3 have been observed in other nanocomposites formed from layered silicates made using very different materials and processes. At the present time, it is not possible to explain the origin of such structures or to know if they are truly representative of the morphology or simply artifacts of the microtoming. It was hoped

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FIGURE 18.3 TEM micrographs of butyl rubber nanocomposites.

that these TEM observations would lead to an experimental assessment of the aspect ratio of the vermiculite particles. However, quantification of this important morphological feature seems essentially impossible because of the difficulties mentioned already and the further complications of defining the lateral dimensions of large particle within the limited field of view dictated by the high magnification needed to observe the particles. From these and other TEM micrographs, one can easily believe that the aspect ratio might be in the range of 100–400; however, we cannot be more specific. 18.1.1.3 Nanoclay Impact on Barrier Properties and Tortuous Path Modeling The tortuous path models and a comparison with both InMat and other nanocomposite data have been reviewed by H.A. Goldberg [10] and Takahashi et al. [9]. These models describe how the permeability of a nanocomposite depends upon the concentration, aspect ratio, dispersion, orientation, and microstructure of the clay platelets in a clay–polymer nanocomposite. Table 18.1 [9,11–14] summarizes the results of the above-mentioned models.

562

Fredrickson and Bicerano model: x ¼ pða=2Þf=lnða=2Þ.

Disc

Gusev and Lusti

a

Disc

Ribbon

Cussler Regular array Random array

Fredrickson and Biceranoa

Ribbon

Filler Type

Nielson

Model

Geometic Details

TABLE 18.1 Detailed Schematics and Formulas for Each Model Discussed in the Text

2

ð1  fÞ ¼ ð1  fÞexp

3:47

h
0:71 i af

 2 x2 ð1  fÞ ¼ 4 1 þ x þ2 0:1245 þx P0 P

P0 P



 2 ð1  fÞ ¼ 1 þ 23 a2 f P0 P

f2

ð1  fÞ ¼ 1  f þ


a2

ð1  fÞ ¼ 1 þ a  f2

P0 P

P0 P

Formulas

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FIGURE 18.4 Large reductions in permeability fit Cussler’s model best.

Nielsen’s model has been applied to clay–polymer nanocomposites for a long time, despite the fact that most materials do not have a ribbon-like microstructure. Many systems that have less than a 10 times reduction in permeability often fit the Nielsen equation quite well. These systems typically have a volume fraction of less than 10% and an aspect ratio of 100 or less. Early work by InMat confirmed the quadratic dependence on the square of the product of aspect ratio and concentration (aw)2 for their system where the aspect ratio and concentration is much higher (Figure 18.4). The uncertainty in the effective aspect ratio is caused by the numerous choices of numerical coefficient that depend upon the detailed microstructure. For very large aspect ratio and concentrations, the Cussler model is difficult to distinguish from the Friederickson and Biceraro equation. 18.1.1.4 Interfacial Interactions and Control Over Flexibility The interfacial interactions between the nanodispersed clay and the polymer matrix are a critical part of the elastomeric barrier coating technology. The water dispersion approach to forming high barrier nanocomposites enables it to control the nanoclays concentration over a much wider range than can be done using conventional compounding and thermal processing. In addition, if there is a strong interaction between the nanoclays and the polymer, the resulting nanocomposite will not remain flexible. This can occur because of the following two effects: . .

A strong interaction will reduce the polymer mobility in the highly confined polymer between the nanoclay plates. If the polymer cannot move over the clay surface, there will be significant strain amplification in the regions at the edges of the clay particles creating sites for mechanical failure.

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FIGURE 18.5 Optical micrographs depicting biaxial strain of elastomeric nanocomposites.

The Tg of the nanocomposite is unchanged from that of the unfilled butyl rubber (
 40 C). This is evidence that there is not a strong nanoclays–polymer interaction in the Air D-Fense products. This is in contrast to the montmorillonite–polyester nanocomposite (Nanolok PT MM) where the polymer phase Tg disappears upon formation of a nanocomposite. The second, and more important result of this control over the interfacial interaction is the fact that its highly filled nanocomposite can still undergo significant strain without damage. Evidence for that is shown in the optical micrographs Figure 18.5. Figure 18.5 shows an earlier version of an elastomeric butyl nanocomposite coating on EPDM rubber on the left. The top left shows cracking of the coating with only 1% strain (the silver dots are used as fiducial marks so that the strain can be measured from the image). At 10–20% strain, the coating is delaminating from the rubber substrate because it cannot undergo the strain. The images on the right, on the other hand, show the commercial Air D-Fense 2000 coating on the same EPDM substrate. The top right figure shows the coating before any strain. The bottom right figure shows the same coating after a biaxial strain of 30% (again this can be determined by the change in spacing of the silver dots). If we look very carefully, we can see some slight change in surface texture at
20–25% strain, but no catastrophic damage. The remarkable flexibility of the Air D-Fense nanocomposite coatings is achieved when the coating is completely dry and tested under desiccating (0% RH) conditions.

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FIGURE 18.6 Helium permeation rate of elastomeric barrier coated natural rubber as a function of percent strain.

This is further evidence that this performance is due to modifications of the clay–polymer interface. The coating on the left that cracks at very low strain will not crack if it is kept at higher (>/¼50% RH) humidity. Since butyl rubber does not change with increasing humidity, the only place a significant amount of water can be absorbed is at the nanoclay–polymer interface. It is clear that this water can reduce the clay–polymer interaction and plasticize the interface. In InMat’s proprietary surface modification technology, it has achieved a similar nanoclay–polymer interface without requiring any moisture or humidity. Figure 18.6 below shows the helium permeation rate through a 0.8 mm natural rubber sheet that was coated with
25 mm of Air D-Fense 2000 as a function of how much the sheet was stretched. One expects the permeation rate to increase with increasing strain even without any damage to the barrier coating because both the substrate and the coating will get thinner. The effect is shown in the top curve that shows the change in helium permeation rate with strain of the uncoated rubber sheet. The important result from this experiment is that even at 40% strain the barrier coating provides a significant improvement in the barrier properties of the rubber sheet, although the permeation rate is higher than expected solely from the change in thickness with strain. The relative (as compared to uncoated, unstrained natural rubber) helium transmission rate for an AD 2000 coated natural rubber substrate is shown. Transmission rates go up as expected from strain thinning. At 40% strain, there is a small increase in the helium permeation rate beyond what is expected from strain thinning. Air D-Fense 2000 and the other water-based elastomeric barrier coating nanocomposites are completely unique in their combination of barrier and compatibility with rubber processing and flexibility. The starting butyl rubber polymer can be used as a coating, but it is neither practical nor cost-effective as a barrier material. That is because one typically would need close to 1 mm in thickness in order to achieve the

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TABLE 18.2 Oxygen Permeability of Elastomeric Barrier Coatings as Compared to Standard Materials Material

O2 Permeability (cc mm/(m2 day atm))

Natural rubber Butyl rubber


2000
100 (70–400 depending upon grade and formulation) 3.0 1.5 0.3

Air D-Fense 2000 (AD 2000) Air D-Fense 3000 (AD 3000) Air D-Fense 5000 (AD 5000)

barrier required in most applications. More commonly, either natural rubber or butyl rubber is used as the barrier material. Table 18.2 shows the oxygen permeability of natural rubber, butyl rubber, and butyl rubber based nanocomposites based on the above discussed technology all normalized to a thickness of 1 mm. 18.1.2 Comparison with Thermally Processed Elastomers Most of the coatings useful in the industry that contain platelet-type fillers are prepared by melt processing, in which solid polymer and solid filler are melted together and mixed at high shear rates. Such melt-processed coatings have 100% solids, and usually use less than about 3% by weight of the platelet fillers. Such coatings do not optimally reduce permeability [15–21].

18.2 APPLICATIONS 18.2.1 Sports Balls and Other Pneumatic Applications Elastomers are often used in products that are pressurized with air or another gas. The elastomer provides both a mechanical support and the container that needs to hold the pressurized gas. The permeation rate through the elastomer determines the stability of the internal pressure. Butyl rubber has an oxygen and nitrogen permeability that is an order of magnitude lower than that of natural rubber. On the other hand, butyl has a mechanical loss tangent that is two to four times higher than natural rubber. It is also more expensive and harder to process than natural rubber. Thus, natural rubber is often used despite its poor ability to maintain pressurized air inside the product. In Table 18.3, we compare the air retention lifetime of several rubber-based pneumatic products. The current technology is what can be achieved with Air D-Fense 2000, while the future potential performance column refers to what can be achieved with Air D-Fense 5000. All values are in months. The pressure change used to calculate the lifetimes was
20% for tennis balls and soccer balls, and 10% for all of the other products.

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TABLE 18.3 Air Retention Lifetime of Pressurized Products Pressure Retention Time (Months)

Product Tennis ball Soccer ball Basketball Mountain bike tire

Current

With Commercial Nanocomposite Coating (Air D-Fense 2000)

With High Performance Nanocomposite Coating (Air D-Fense 5000)

0.4 0.3 6.7 0.5

1.4 3.5 11.7 0.7

8.1 24.9 44.4 1.8

It is important to note that all the pressure retention lifetimes shown above could be improved by another factor of two if 100% nitrogen gas were used instead of air. In pneumatic products, the steady-state permeation rate and pressure loss rate are the important performance parameters that determine if a barrier coating is valuable. 18.2.1.1 Tennis Balls Tennis balls are manufactured from natural rubber because of its low mechanical loss, which leads to its relatively high bounce. A butyl rubber ball would hold air pressure better, but it would not bounce very well. The elastomeric barrier coatings described herein have been the only material solution to the problem of pressure retention in tennis balls that has been commercialized. Improvements in pressure retention have been achieved by changing the gas composition in the tennis ball. The use of 100% nitrogen will improve the pressure retention by approximately two times. This would work for all pneumatic products. A typical tennis ball will lose its air pressure within a few days after the can is opened. In fact, standard pressurized tennis balls are sold in a pressurized can. There have been several attempts to solve this problem of air pressure loss. An early effort to replace the pressurizing gas with sulfur hexafluoride (SF6) failed [22]. Two reasons that have been cited include the change in the sound the ball makes during play due to the different density and acoustic properties of the heavier gas, and the difficulty in quality control when gas mixtures are used to fill the tennis balls. It was important to fill the ball with an air–SF6 mixture because if the ball was filled only with SF6 (which permeates very slowly through the rubber) air would slowly permeate into the ball increasing its pressure. Several patents have been filed addressing these issues [23–26]. The idea of using a barrier coating inside the ball was also not new. The problem was that one needed a sufficiently good gas barrier so that a thin coating could be used and the weight and bounce of the ball would not be significantly changed. Until the elastomeric barrier coating technology was developed, the only other barrier materials were not compatible with the rubber manufacturing process used to make tennis balls. During that process, a coating has to withstand greater than 100 C temperatures during curing, as well as a 5% change in volume when the ball is pressurized. That is, before it is covered with felt and placed in the can.

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The elastomeric barrier coating technology is uniquely suited to solve these problems. A thin (20–30 mm) coating provides a barrier equivalent to 10 mm of natural rubber or three to four times that of the 3 mm thick tennis ball. It is designed to go through a standard rubber vulcanization process, which begins at about 100 C, and can be carried out at temperatures up to 170 C. The thin coating adds less than 0.5 g of weight to a tennis ball. This is much less than the standard deviation of the weight of individual balls, and is less than 5% of the total weight. These properties led Wilson Sporting Goods to commercialize its Double Core Tennis Ball in 2001 using the elastomeric nanocomposite barrier coating technology. This was the first commercial application of a nanocomposite barrier coating, and as such it received significant publicity [27]. 18.2.1.2 Soccer Balls Soccer balls typically use a natural rubber bladder. Butyl rubber bladders are also used as they hold the air pressure better, but they feel stiffer. All IFF (International Football Federation) approved balls use a natural rubber bladder. The elastomeric barrier coating technology described herein coated on a natural rubber bladder can provide the air retention of a butyl bladder while not significantly altering the feel and bounce of a ball made with a natural rubber bladder. Figure 18.7 shows the percent pressure loss of soccer balls made with natural rubber bladders both with and without an Air D-Fense 2000 coating. It clearly shows the large increase in time between reinflations in order to stay in the playable range (80–90% of ideal inflation). It also shows that after several months, there is still significant internal pressure. This makes balls with an Air D-Fense 2000 coating look much better when left on the shelf in retail stores for extended periods of time. The process of applying the elastomeric barrier coating technology on a soccer ball has also been tested extensively. Natural rubber bladders are currently produced by a dip coating process whereby a form is dipped into a suspension of natural rubber

FIGURE 18.7 Pressure loss of coated and uncoated soccer ball bladders.

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latex right after being coated with a coagulating agent. The thickness of the bladder is controlled by the dipping time. After removing the bladder from the natural rubber suspension, the bladder is dried, cured, and washed. In order to use essentially the same processing equipment, a dip coating process was developed whereby the washed bladders could be dipped in the elastomeric barrier coating suspension, dried, and cured a little more. The coating thickness of the nanocomposite barrier is determined by its solid content and rheology. Since the bladders were typically turned inside out when they were removed from the dip coating forms and washed, the barrier coating was on the inside. After the barrier coating was dried and cured, the bladder would be turned inside out again when the coated bladder was removed from the InMat designed bladder holder. Thus, the bladder would have a barrier coating on its inside. The bladder was then placed in a fixture whereby the inflation valve was glued in place. The importance of improved pressure retention in soccer balls is clear from the fact that a large (
40%) section of the soccer ball market uses a butyl rubber bladder instead of a natural rubber bladder. These bladders cost more than natural rubber bladders, and change the bounce and feel of the ball. Manufacturers have to modify the design of the other soccer ball layers to compensate for the changes caused by butyl bladders. More recently, Primosports has introduced an alternative bladder technology to address this issue [28]. Although nanocomposite barrier coatings in soccer balls have not been commercialized, we believe there is still a strong potential for this technology in this application. 18.2.1.3 Basketballs Unlike tennis balls and soccer balls, the amount of rebound (or bounce height) required in a basketball is not that high. Therefore, most basketballs already use butyl rubber in their construction. Therefore, the basketballs hold their air pressure much better than soccer balls and tennis balls. On the other hand, basketball players are much more sensitive to small pressure changes in a ball, as it is constantly being bounced by the player. Thus, even small changes in pressure over time are an issue with the consumer. The 6.7 month time it takes a basketball to lose 10% of its pressure was considered enough of a problem, that Spalding, the leading producer of basketballs, has introduced at least two products over the years to solve this problem [29]. In fact, Spalding claims that its standard balls only maintain their bounce for about 1 month. Part of the issue with basketballs is that the butyl rubber used to produce them can be made from a wide variety of compositions. Typically the butyl rubber is blended with natural rubber to decrease cost and improve its processability. This also leads to large variations in its gas permeability. The use of a nanocomposite coating to reduce the air permeation of the ball would resolve this problem while making the air pressure retention much less sensitive to the rubber composition used. 18.2.1.4 Other Inflated Sports Balls There are several other sports balls that use an internal bladder inflated with air. They include American footballs, rugby balls, and volleyballs. All have very much the same issues as discussed above for soccer balls. In particular, natural rubber, butyl rubber, and blends of the two are the

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most commonly used materials. American footballs have used a variety of polymers for the air bladder. Apparently, the elasticity of the bladder is often considered less important than its air pressure retention, as relatively rigid materials such as polyesters and nylon have even been used. 18.2.1.5 Bicycle and Specialty Tires and Tubes Most bicycle tires use butyl rubber tubes to maintain air pressure. The thickness of the butyl rubber determines the permeation rate and thus the rate at which pressure is lost. The tube adds a significant amount to the weight of the tire and to the rolling resistance. Weight is one of the most important performance factors in bicycle tubes. Some bicycle racers use a natural rubber tube, which is very low weight and has lower rolling resistance. As discussed above, the natural rubber has much lower loss tangent than butyl rubber, and thus much less energy is lost when it is deformed. Natural rubber tubes are rarely used despite this fact because of their high pressure loss rate. Other factors including puncture resistance and failure mechanisms when the tube is punctured are also very important. In addition to butyl rubber and natural rubber, other elastomers are often used as part of the composition of bicycle tire tubes. EPDM rubber is easier to process and is often blended with butyl rubber. On the other hand, it is more permeable than natural rubber. Polyurethane elastomers could also be considered for tubes and could be made into very lightweight structures. The use of a nanocomposite barrier coating on any of these materials would separate the important considerations of weight, and puncture resistance from air pressure retention, and thus could lead to a new class of low weight, high performance bicycle tubes. 18.2.1.6 Pressure Accumulators Rubber bladders are often used to store mechanical energy through the compression of the gas inside. This includes industrial equipment such as molding machines, as well as recently commercialized mechanical hybrid vehicles in which the energy from breaking is recovered by compressing a rubber bladder. In these devices, the bladder is initially pressurized, and the energy storage mechanism is through compression. Performance depends upon maintaining the pressure inside the bladder. Thus, the permeation rate through the bladder is an issue, and the use of a nanocomposite barrier coating is being considered. 18.2.1.7 Hoses Rubber hoses are often used to transfer chemicals and gases. Permeation through the hose is often an issue, and the use of a nanocomposite barrier coatings has been considered. Automotive hoses for fuel and refrigerants are one of the largest areas where hose permeation is critical. The emission standards for both hydrocarbons and refrigerants for automobiles have in fact led to the reduced use of elastomers for these hoses. Many of the automotive applications that use rubber hoses now use more rigid assemblies made from less permeable thermoplastics. In places where a high flexibility rubber hose is still required, the use of an elastomeric nanocomposite barrier coating should be considered. The same issues apply to nonautomotive uses of rubber hoses for chemical and gas transfer.

APPLICATIONS

571

18.2.2 Breakthrough Time Applications All of the above applications make use of the low steady-state permeability of clay–polymer nanocomposite coatings. This means that the short time behavior after the product is initially pressurized is not important. This is not typically the situation in chemical protective products. In those products, what is important is how long it takes for the first small amount of material to come through the protecting layers. The breakthrough time is approximately proportional to the thickness squared and inversely proportional to the diffusion constant for a uniform slab of a single material. Thus, thin films have a very short breakthrough time. When a very thin barrier coating is put on a very permeable substrate, the steady-state permeation rate can be changed significantly, while the changes in breakthrough time can be much smaller. The steady-state permeation rate does not depend on how a thin barrier material is deployed on a more permeable substrate. This means that if the total thickness of the barrier material and substrate is the same, the steady-state permeation rate will be the same for samples with a single barrier layer, or the barrier layer split into two or more coatings and separated by some or all of the more permeable substrate. This is not the case for the breakthrough time. The breakthrough time can depend critically on the way the barrier material is deployed. In particular, if half the barrier layer is put on one side of the more permeable substrate, and the other half is put on the other side, there can be much larger increases in breakthrough time when compared with putting all the barrier material on one side. 18.2.2.1 Chemical Protective Gear The most common example of this kind of product is chemical protective gloves and other chemical protective gear. What is important is how long it takes for a hazardous amount of a toxic chemical to permeate. The long-term rate is not usually an issue. Many of these products are made from elastomers so that they have adequate flexibility. Given the high permeability of elastomers, the protective products are often relatively thick. Thus, a surgeon’s glove that only is designed to block microorganisms can be made from very thin natural rubber latex. These gloves are typically penetrated by hazardous chemicals in seconds, and thus do not provide any chemical protection. For chemical protection, butyl rubber gloves that are 12–35 mils thick are often used. They can provide 12–24 h of protection from mustard gas, for example. The protection time (often called the breakthrough time) is determined by how long it takes for
5 mm/cm2 of mustard gas to penetrate the glove, and increases approximately with the square of the glove thickness [30]. When chemical protection products were initially made from multiple materials, it was usually found that the breakthrough time correlated well with the steady-state permeation rates. In these products, the different layers typically did not differ from each other by more than a factor of 10.

572

NONAUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES

FIGURE 18.8 Chemical warfare agent barrier performance (MIL-STD-282).

When work began using nanocomposite barrier coatings in combination with much more permeable elastomers, it was found that large improvements could be achieved in the breakthrough time. The first example is shown below in Figure 18.8, where a very permeable natural rubber sheet was coated on both sides with a thin nanocomposite barrier coating, and was able to achieve the same protection from mustard gas as a thick butyl rubber glove. The nanocomposite coating was Air D-Fense 2000, and only 20 mm was coated on each side. As already mentioned, the fact that such thin coatings could make such a large change in breakthrough times was initially surprising. It was very quickly discovered that when a very impermeable coating is placed on a much more permeable substrate, the total thickness of the impermeable layers was not the only important design variable. It also was important to consider whether the barrier layers were deployed as a single layer or as multiple layers. What was discovered was that if two (or more) barrier layers are used, the thickness of the separating more permeable layer was of critical importance. In addition, much larger reduction in breakthrough time could be achieved with this type of construction [31,32]. A significant amount of work was done under a government SBIR grant to demonstrate the utility of the elastomeric nanocomposite barrier coating technology for protection of soldiers from chemical warfare agents [33]. 18.2.2.2 Reduction of Blooming and Staining Single layers of nanocomposite barrier coatings can also provide significant improvements. Because of process cost concerns, they are sometimes used despite the clear advantage of using at least two layers. Recently, an application in which staining of rubber processing chemicals and oils needed to be delayed was commercialized. InMat’s elastomeric nanocomposite barrier coating technology was used to increase the time before these chemicals caused a color change from a few days to about 1 month.

REFERENCES

573

REFERENCES 1. Maxfield, M.; Christiani, B. R.; Sastri, V. R. U.S. Patent 5,514,734 (to AlliedSignal Inc.) (1996). 2. Christiani, B. R.; Maxfield, M. U.S. Patent 5,5,747,560 (to AlliedSignal Inc.) (1998). 3. Maxfield, M.; Christiani, B. R. WO 93/11190 (to AlliedSignal Inc.), (1993). 4. Feeney, C. A.; Farrell, M.; Tannert, K.; Goldberg, H. A.; Lu, M.; Grah, M. D.; Steiner, W. G.; Winston, P. B. U.S. Patent 6,087,016 (to InMat LLC) (2000). 5. Feeney, C. A.; Farrell, M.; Tannert, K.; Goldberg, H. A.; Lu, M.; Grah, M. D.; Steiner, W. G.; Winston, P. B. U.S. Patent 6,232,389 (to InMat LLC, Michelin Recherche et Technique S.A.) (2001). 6. Feeney, C. A.; Goldberg, H. A.; Farrell, M.; Karim, D. P.; Oree, K. R. U.S. Patent 7,078,453 (to InMat Inc.) (2006). 7. Feeney, C. A.; Goldberg, H. A.; Farrell, M.; Karim, D. P.; Oree, K. R. U.S. Patent 7,119,138 (to InMat Inc.) (2006). 8. Goldberg, H. A., et al. Materials Research Society Symposium Proceeding, 733E:P. T.4.7.1–4.7.6 (2002). 9. Takahashi, S.; Goldberg, H. A.; Feeney, C. A.; Karim, D. P.; Farrell, M.; O’Leary, K.; Paul, D. R. Gas barrier properties of butyl rubber/vermiculite nanocomposite coatings. Polymer, 47:3083–3093 (2006). 10. Goldberg, H. A. NICHE Presentation, June 2004. 11. Nielsen, L. E. Models for the permeability of filled polymer systems. J. Macromol. Sci. (Chem.), T. AI:929–942 (1967). 12. Lape, N. K.; Nuxoll, E. E.; Cussler, E. L. Polydispersed flakes in barrier films. J. Membr. Sci., T. 236:29–37 (2004). 13. Gusev, A. A.; Lusti, H. R. Rational design of nanocomposites for barrier applications. Adv. Mater., T. 13:1641–1643 (2001). 14. Fredrickson, G. H.; Bicerano, J. Barrier properties of oriented disk composites. J. Chem. Phys., T. 110:2181–2188 (1999). 15. Mowdood, S. K.; Kansupada, B. K. U.S. Patent 4,857,397 (to The Goodyear Tire & Rubber Company) (1989). 16. Elpass, C. W.; Kresge, E. N.; Peiffer, D. G.; Hseih, D.-T.; Chludzinski, J. J. WO 97/00910 (to Exxon Research and Engineering Co.) (1997). 17. van Amerogen, G. J. Diffusion in elastomers. Rubber Chem. Tech., T. 37:1065–1152 (1964). 18. Cussler, E. L. et al. Barrier membranes. J. Membr. Sci., T. 38:161–174 (1988). 19. Nielsen, L. E. Models for the permeability of filled polymer systems. J. Macromol. Sci. Chem., T. A1:929 (1967). 20. Bharadwaj, R. K. Modeling the barrier properties of polymer-layered silicate nanocomposites. Macromolecules, T. 34:9189 (2001). 21. Fredrickson, G. H.; Bicerano, J. Barrier properties of oriented disk composites. J. Chem. Phys., T. 110:2181 (1999). 22. Koziol, D. L.; Reed, T. F. U.S. Patent 4,098,504 (to The General Tire & Rubber Company) (1978). 23. Reed, T. F. U.S. Patent 4,166,484 (to The General Tire & Rubber Company) (1979).

574

NONAUTOMOTIVE APPLICATIONS OF RUBBER–CLAY NANOCOMPOSITES

24. Reed, T. F. U.S. Patent 4,248,275 (to The General Tire & Rubber Company) (1981). 25. Reed, T. F.; Ritzert, R. K. U.S. Patent 4,300,767 (to The General Tire & Rubber Company) (1981). 26. Papinsick, J.; Oransky, J. J.; Sircar, S. U.S. Patent 4,358,111 (to Air Products and Chemicals,; Inc.) (1982). 27. Time Magazine. September 23, 2002, A14. 28. www.primosport.com. 29. http://www.beststuff.com/categories/fromthewire/spalding-solves-inflation-issue-newneverflat-basketball.html. 30. Lindsay, R. S.; Longworth, T. L.; Johnson, M. A.; Baranoski, J. M.; Hannigan, J. B. Domestic Preparedness Program, Liquid Sulfur Mustard and Sarin Challenge/Vapor Penetration Swatch Testing of Glove Set, Chemical Protective NSN: 8415-01-033-3517; report # ERDEC-TR-536; October 1998. 31. Karim, D. P.; Goldberg, H. A.; Feeney, C. A.; Farrell, M. US Publication U.S. 2006/ 0110615 A1 (to InMat Inc.) (2006). 32. 2004 Scientific Conference on Chemical and Biological Defense Research, Hunt Valley, MD on November 15–17, 2004. 33. SBIR Contract Number: DAAD16-03-C-0041 Final Report.

INDEX

Abrasion, apolar diene rubber nanocomposites, 392–393, 398–399 Accelerators: ethylene-propylene-diene-monomer rubber, in situ intercalation preparation, 469–474 rubber-clay nanocomposite vulcanization and permeability, 356–358 sulfur vulcanization, 278–282 vulcanization reaction, organo-clay kinetics, 284–289 Acid-base equilibria, rubber-clay nanocomposites, vulcanization activation, 139–141 Acrylated oleic methylester (AOME), solution and in situ polymerization, 170 Acrylonitrile butadiene rubber (NBR)-clay nanocomposite, rheology, 250–253 Additives: latex compounding, natural butadiene rubber, 412

thermoplastic polyurethane nanocomposites, thermal properties, 504–507 Adhesion properties, butyl rubber nanocomposites, 456–457 Aggregation: clays and clay minerals, 6 pristine clays, rubber matrix, 189–190, 194–197 Aging, automotive rubber-clay nanocomposites, 539–541 Air barrier properties: automotive rubber-clay nanocomposites, 538–539 butyl rubber clay nanocomposites, 449–452 Air-D-Fense products, 564–566 Air oven aging, bromobutyl rubber, exxpro, and nanocomposite innerliner compounds, 355–356 Alkenylamines, rubber-pristine clay composites, 197–206

Rubber-Clay Nanocomposites: Science, Technology, and Applications, First Edition. Edited by Maurizio Galimberti. Ó 2011 John Wiley & Sons, Inc. Published 2011 by John Wiley & Sons, Inc.

575

576

INDEX

Alkenyl groups, ammonium cation substituents, 215 Alkylammonium: chain structure and dynamics, 101–111 dynamics and diffusion, 110 molecular simulation, 111 packing density and self-assembly, 102–110 clay cationic exchange capacity, 25–27, 50–58 clay mechanical properties, 121–123 organophilic clays, 48–58 polymer layer separation and miscibility, 115–121 cleavage energy, 116–121 exfoliation, thermodynamic polymer matrix model, 115–116 surface energy, 121 thermal properties, 111–115 reversible melting transitions, interlayer alkyl chains, 111–113 solvent evaporation and thermal elimination, 113–114 Alkylphosphonium, organophilic clays, cation exchange, 50–58 Alkyl surfactants, solvent evaporation and thermal elimination, 114–115 Allophanes: availability, 33–35 TO criteria, 16 Allylamine, latex compounding, rubber-clay nanocomposites, 155–157 Amines, thermal decomposition mechanism, 135 3-Aminopropyltriethoxysilane, dispersion, 386–387 Ammonium cations: BIMSM ionomers, 442–443 organophilic clays, rubber vulcanization, 288–289 organophilic clays, cation exchange, 50–58 rubber-clay nanocomposites, 206–212 hydrogenated tallow and benzyl groups, 213–215 long- and short-chain alkenyl groups, 215 long-chain alkenyl substituents, montmorillonite modification, 219–228

montmorillonite and bentonite, 207–212 polar montmorillonite, 215–219 Ammonium cations, tallow group composites, 220–228 Amphophilicity, organophilic clay minerals, 46 Amphoteric surface chemistry: clay minerals, 24 clay rheology, 32–33 Anionic clay minerals: layered structure, 21–22 rubber-clay nanocomposites, vulcanization activation, 139–141 Anionic exchange capacity (AEC), clay minerals, 24–27 Apolar diene rubber composites: applications and future research, 398–399 barrier, 393–396 basic properties, 369–371 characterization, 383–388 cure characteristics, 377–379 dispersion detection, 380–383 fire resistance, 396–397 friction/wear/abrasion, 392–393 mechanical properties, 388–392 plate structure, 397–398 preparation, 371–377 latex compound, 371–373 melt blending, 374–377 solution mixing, 373–374 Arrhenius equation: rubber vulcanization, curing mechanics, 291–297 thermoplastic polyurethane nanocomposites, decomposition kinetics, 499–501 Arrhenius-Frenkel-Eyring equation, pristine clay modification, rubber-clay nanocomposites, 256–257 Aspect ratio: butyl rubber-clay nanocomposites, 436–437 clay aggregation and dispersion, 6 elastomer nanocomposite gas permeability, 348–349 organophilic clays, 47–49 pristine silicates, 191

INDEX

rheology of rubber-clay nanocomposites, 245–246 water-based nanocomposites, barrier mechanisms, 558–561 Atomic force microscopy (AFM), apolar diene rubber nanocomposites, 380–383 Automotive applications, rubber-clay nanocomposites: aging and ozone resistance, 539–541 barrier/air retention properties, 538–539 basic principles, 525–526 disadvantages, 548–549 elastomeric components, 531 elastomeric nanocomposites, 531–548 future research issues, 548–550 organoclay minerals, 534 polyurethane-organoclay composites, 544–545 processability, 542–544 rubber, 526–531 rubber-clay vs. other nanofillers, 534 solvent resistance, 541–542 tire nanocomposites, 545–548 weight and balanced mechanical properties, 535–538 Backside temperature, thermoplastic polyurethane nanocomposite flammability, 508 Balanced mechanical property, automotive rubber-clay nanocomposites, 535–538 Barrett-Joyner-Halenda (BJH) method, clay surface area and porosity, 22–24 Barrier mechanisms: apolar diene rubber nanocomposites, 393–396 automotive rubber-clay nanocomposites, 538–539 nitrile rubber composites, 423–424 water-based nanocomposites, 557–566 dispersion and orientation, 558–561 interfacial interactions and flexibility control, 563–565 tortuous path modeling, 561–563 Basal plane spacing, alkylammonium chains, packing density and selfassembly, 107–110

577

Basketballs, water-based nanocomposites, 569 Belts (automotive), rubber applications, 529 “Bentones,” organophilic clays as, 48–49 Bentonite: availability, 33–35 commercial applications, 88–89 diffusion, 87–89 geological occurrence, 89 industrial treatment, 87–90 mining, 89–90 modification, 90–91 processing, 90–93 rubber-clay nanocomposites: ammonium cations, 207–212 pristine clay modifiers, 190–191 rubber-pristine clay composites, 193–194, 195 primary alkenylamines, 198–202 structure, 87–89 swelling properties, 17, 20–21 Benzyl groups, ammonium cations, 213–215 Bicycle tires, water-based nanocomposites, 570 Bilayers, organophilic clays, cation exchange, 54–58 Bis(dialkldithiocarbamate)zinc(II) (ZDAC), 469 Blooming, water-based nanocomposites, 572 Bragg diffraction model, rubber-clay nanocomposite morphology, 182 Brake hose, rubber applications, 528 Brake vacuum hose, rubber applications, 527 Breakthrough time applications, water-based nanocomposites, 571–572 Brominated isobutylene-co-para-methylstyrene (BIMSM) rubber-clay nanocomposites: butyl rubber, 438 air permeability, 452 processability, 457 reinforcement properties, 452–454 solution method, 446–447 vulcanization, 454–456 innerliner formulations, 353–356 ionomers, 439, 441–443

578

INDEX

Brominated (Continued ) montmorillonite nanocomposite thermodynamics, 358–362 permeability patterns, 349–350 pristine clay modification, 202 rheology, 257 Bromobutyl screening compound, permeability measurement, 346–347, 353–356 Bulky inorganic oligomers, intercalation, 28 Butadiene-isoprene-isobutadiene rubber (BIIR) nanocompound, melt compounding, 163–169 Butadiene rubber (BR): apolar diene rubber nanocomposites: dispersion characterization, 385–387 mechanical properties, 390–392, 393 melt blending, 375–377 diene rubber composites, 222–225 melt compounding, 159–160 polar group ammonium cation modification, 217 pristine clay modification, rubber-clay nanocomposites, 190–191 primary alkenylamines, 200–202 rubber-clay composites, 195 ammonium cations, 207, 213 rheology, 247–250 Butadiene-styrene-vinylpyridine rubber, clay interaction/affinity with, 153–157 Butyl rubbers: copolymerization, 369–371 ionomers, 439–443 key properties and applications, 431–433 low molecular mass substances, 444 maleic anhydride-grafted polymers, 443–444 morphology, 433–435 permeability, 343–345 rubber-clay nanocomposites, 433–435 adhesion properties, 456–457 air barrier properties, 449–452 aspect ratio talc filters, 436–437 compatibilizer systems, 438–444 hydrotalcite clays, 435–436 mica clays, 437 montmorillonite clays, 435 preparation methods, 444–448 processability, 457

reinforcement properties, 452–454 sepiolite, 438 vermiculites, 437–438 vulcanization properties, 454–456 solution-based preparation, 374 Calcium silicate hydrates, layered structure, 21 Carbon black: apolar diene rubber nanocomposite mechanics, 387–392 automotive nanocomposites, aging and ozone resistance, 539–541 brominated isobutylene-co-para-methylstyrene-montmorillonite nanocomposite thermodynamics, 362 reinforced elastomeric materials, 328–333 Carbon nanotubes: automotive applications, 531 fire resistance, 424–425 Carboxylic pendant groups, natural butadiene rubber (XNBR): basic properties, 409–410 demand loading conditions, 420–421 melt compounding, 413–414 polarization effects, 425 Carboxyl-terminated butadiene (CTB): apolar diene rubber nanocomposites: mechanical properties, 391–392 melt blending, 375–379 butadiene-rubber nanocomposite rheology, 250 Carboxyl-terminated butadiene acrylonitrile (CTBN) copolymer, apolar diene rubber nanocomposites, melt blending, 376–380 Carreau model, polyepichlorohydrin rubberclay nanocomposite rheology, 259–261 Cation exchange capacity (CEC): alkylammonium chains, 101–111 cleavage energy, 116–121 clay minerals, 24–27 purification, 93–97 organophilic clays: dispersion maximization, 69–70 rubber vulcanization, 288–289 solution-based, 49–58

INDEX

Cationic clays: layered silicates, 21 layer organization, 8–11 nonphyllosilicate layers, 21 rubber-clay nanocomposite morphology, 182–183, 187 Cellulose whiskers, automotive applications, 531 Charge density, clay cationic exchange capacity, 25–27 Chemical protectiive gear, water-based nanocomposites, 571–572 Chemical treatments: organophilic clays, 71 pristine clay modification, 191 Chlorites, TOT criteria, 14 Chlorobutyl rubber: apolar diene rubber nanocomposites, 392 permeability, 344–345 Chlorosulfonated polyethylene (CSM), rubber applications, 527 Clay-polymer nanocomposites (CPN): as fillers, 35–37 kaolin intercalation, 28–29 one-pot formation, 66 organophilic clays, 47–49, 72–75 solid-state intercalation, 58–59 swelling properties, 20, 31 synthetic clay availability, 35 Clays and clay minerals: automotive nanofillers, 534 availability, 33–35 basic properties, 3–5 chemical formula, 18 as fillers, 35–37 layer organization, 8–22 allophane and imogolite group, 16–17 anionic minerals, 21–22 cationic layered silicates, 21 cationic/neutral clay materials, 8–11 chlorites, 14 interstratified minerals, 14–15 kaolinite and serpentine, 11 micas, flexible and brittle, 14 nanostructures, 20–21 nonphyllosilicate cationic layered minerals, 21 pyrophyllite and talc, 14 sepiolite and palygorskite group, 15–16

579

smectites, 14 swelling clays, 17, 20–21 TO and TOT groups, 11–17 vermiculites, 14 mechanical properties, 121–123 melt compounding, rubber-clay nanocomposites, microstructure evolution, 162–163 multiscale organization, 6–8 delamination/exfoliation vs. stacking, 6–8 dispersion vs. aggregation, 6 nanometric architecture, 3–4 natural clays, availability, 33, 35 organic substance reactions, 97–98 organophilic minerals: chemical treatments, 71 clay-polymer nanocomposites, 72–75 dispersion maximization, 69–70 hydrophilic/lipophilic balance, 45–47 nanomaterials, 66–69 physical treatments, 71–72 polymer technology, 47–49 synthesis procedures, 49–66 thermal stability, 70–71 particle size modification, 99 physicochemical properties, 22–33 cation/anion exchange capacity, 24–27 interlayer intercalation and confinement, 27–30 rheology, 31–33 surface area and porosity, 22–24 surface chemistry, 24 swelling, 30–31 purification, 93–97 clay concentration, 94 swelling time, 94–95 temperature, 95–97 synthetic clays, availability, 33–35 Clay stacking, rubber-clay nanocomposite formation, 228–232 Clay-surfactant interface, alkylammonium chains, diffusion dynamics, 110 Cleavage energy, polymer layer separation and miscibility, 116–121 Cloisite organoclay, 64–65 thermoplastic polyurethane nanocomposites, 513–516

580

INDEX

Coagulating agents, rubber-pristine clay composites, 192–194 emulsion compounds, 196–197 Coatings, reinforced elastomeric materials, 330–333 Co-coagulation methods, latex compounding, rubber-clay nanocomposites, 149–153 Cocondensation, grafted organoclays, 62–64 Coefficient of friction (COF), nitrile rubber composites, 421–422 Coefficient of thermal expansion (CTE), thermoplastic polyurethane nanocomposites, 512–513 Collapse, melt compounding, rubber-clay nanocomposites, 163 Compatibilizers: butyl rubber nanocomposites, 438–440 melt compounding, rubber-clay nanocomposites, 161–162 olefinic nanocomposites, 467–469 reinforced elastomeric materials, 329–333 Compression set determination, reinforced elastomeric materials, 317–319 Cone calorimetry: thermoplastic polyurethane nanocomposite flammability, 508–509 thermoplastic polyurethane nanocomposites, 496 Confinement: clay interlayer space, 27–30 melt compounding, rubber-clay nanocomposites, microstructure evolution, 163 Consistency index, fluoroelastomer-clay nanocomposite rheology, 254–257 Continuum-based micromechanical models, reinforced elastomeric materials, 334–335 Cost issues, automotive elastomers, 548–549 Coulomb interactions, alkylammonium chains, cleavage energy, 119–121 Coupling agents, melt compounding, rubberclay nanocomposites, 161–162 Crack propagation curves, reinforced elastomeric materials, cyclic loading, fracture mechanics testing, 327–328

Creep properties, ethylene-propylene monomer nanocomposites, 482–483 Cross-link density, reinforced elastomeric materials, 332–333 Cross-linking reactions: melt compounding, rubber-nanoclay composites, 166–169 vulcanization effects, 169–170 reinforced elastomeric materials, 330–333 rubber vulcanization, 278–283 peroxide, 282–283 sulfur, 278–282 sulfur vulcanization, 279–282 Crystallization: apolar diene rubber nanocomposites, 389–392 reinforced elastomeric materials, 332–333 Crystallographic molecular structures, organophilic clays, cation exchange, 54–58 Curatives: automotive elastomers, 544 melt compounding, rubber-clay nanocomposites, 162 Cure rate index (CRI), rubber vulcanization, curing mechanics, 278 Curie point, thermoplastic elastomers, 495 Curing bladders, butyl rubber reinforcement properties, 453–454 Curing kinetics: apolar diene rubber nanocomposites, 377–379 nanocomposite vulcanization and permeability, 356–358 rubber vulcanization: mechanics, 276–278 nitrile composites, 414–416 organoclay activation, 137–141, 283–289 reaction mechanics, 276–278 rubber cross-linking systems, 278–283 Cyclic loading, reinforced elastomeric materials, fracture mechanics, 326–328 Cyclohexylbenzothiazole sulfenamide (CBS), vulcanization and permeability, 356–358

INDEX

Decomposition kinetics, thermoplastic polyurethane nanocomposites, 497–501, 498–501 Deformability, reinforced elastomeric materials, tensile testing, 310–313 Deintercalation, melt compounding, rubberclay nanocomposites, 163 Delamination, clay aggregation and dispersion, 6–8 Delta torque, rubber vulcanization, curing mechanics, 278 Demanding loading conditions, nitrile rubber composites, 420–421 Density functional theory (DFT): alkylammonium chains, molecular simulation, 111 clay mineral mechanics, 121–123 Density properties, thermoplastic polyurethane nanocomposites, 497 DesmopanÔ, 492–493 dispersion, 502–503 thermal properties, 504–507 Diaphragms (automotive), rubber applications, 529 Dibutylthiourea (DBTU): brominated isobutylene-co-para-methylstyrene-montmorillonite nanocomposite thermodynamics, 362 nanocomposite vulcanization and permeability, 356–358 vulcanization, 455–456 Diene rubber composites: apolar compounds: applications and future research, 398–399 barrier, 393–396 basic properties, 369–371 characterization, 383–388 cure characteristics, 377–379 dispersion detection, 380–383 fire resistance, 396–397 friction/wear/abrasion, 392–393 mechanical properties, 388–392 plate structure, 397–398 preparation, 371–377 polar group ammonium cation modification, 217–219 tallow ammonium cations, 220–228

581

Differential scanning calorimetry: rubber vulcanization, curing mechanics, 290–297 thermoplastic polyurethane nanocomposites, 497 Diffusion coefficient: butyl rubber elastomers, 345 ethylene-propylene-diene-monomer rubber gas permeability, 473–474 Diffusion dynamics, alkylammonium chains, clay-surfactant interface, 110 Dioctadecylammonium, reversible melting transitions, 113–114 Dioctahedral smectites, swelling properties, 17, 19–20 Direct synthesis, grafted organoclays, 62–64 Dispersion: apolar diene rubber nanocomposites, 379–387 characterization, 383–387 detection, 380–383 clays and clay minerals, 6 nitrile rubber-clay composites, 410–414 solution-base preparations, 410–411 organophilic clays, maximization, 69–70 rubber-clay nanocomposite morphology, 184 pristine clays, 187–192 thermoplastic polyurethane nanocomposites, 501–503 water-based nanocomposites, barrier mechanisms, 558–561 Distribution, rubber-clay nanocomposite morphology, 184 pristine clays, 187–192 Dithiocarbamate vulcanization accelerator, melt compounding, rubber-clay nanocomposites, 164–169 DLVO (Derjaguin-Landau-Verwey-Overbeek) theory, clay rheology, 31–33 Double-edge-notched tension (DENT) specimens, reinforced elastomeric materials, 323–326 Douglas-Garbochi approximation, rheology of rubber-clay nanocomposites, 245–246 d-spacing, elastomer nanocomposite permeabilitgy, 351–352

582

INDEX

Dynamic-mechanical (thermal) analysis (DM(T)A): nitrile rubber composites, 418–421 reinforced elastomeric materials, 307–310 Dynamic storage moduli, ethylene-propylene monomer nanocomposites, temperature dependence, 481–482 Edge-group silanization, organoclay grafting, 61 Elastomers. See Rubber-clay nanocomposites Electrolyte flocculants, latex compounding, rubber-clay nanocomposites, 149–153 Electron microscopy, water-based nanocomposites, orientation, 559–561 Emulsion compounds, rubber-pristine clay composites, 196–197 Entropy-elastic range: brominated isobutylene-co-paramethylstyrene-montmorillonite nanocomposite thermodynamics, 360–362 reinforced elastomeric materials, dynamicmechanical (thermal) analysis, 308–310 Epichlorohydrin rubber (ECO), melt compounding, 159–160 Epoxidized natural rubber (ENR): apolar diene rubber nanocomposites: dispersion detection, 381–383 melt blending, 376–377 solution-based preparation, 374 automotive elastomers, 543–544 diene rubber composites, 222 melt compounding, 159–160, 414 reinforced elastomeric materials, 329–333 rheology, 243–246 Ethylene glycol monoethyl ether (EGME) adsorption, clay surface area and porosity, 23–24 Ethylene-natural rubber (ENR), rubber-pristine clay composites, 195 Ethylene-propylene-diene-monomer (EPDM) rubber: ammonium cations, 209–212

automotive rubber-clay nanocomposites, 536–538 basic properties, 465–466 chemical factors, 163–166 compatibilizers, 467–469 diene rubber composites, 225–228 gas barrier properties, 473–474 Kevlar-filled, kinetics, 513–514 mechanical properties, 476–479 melt compounding, 159–160 microstructure evolution, 163–169 mineral compounds, 466–467 notch sensitivity, 317 physical factors, 166–169 pristine clay modification, rubber-clay nanocomposites, 191 primary alkenylamines, 202 reinforced elastomeric materials, fracture mechanics, 322–323 rheology, 254, 474 rubber-pristine clay composites, 195 in situ intercalation preparation, 469–473 stability, 475 vulcanization reaction, organo-clay kinetics, 285–289 water-based nanocomposites, 564–566 Ethylene-propylene monomer (EPM)-clay nanocomposites: basic properties, 465–466 compatibilizer systems, 467–469 creep properties, 482–483 mineral compounds, 466–467 preparation and characteristics, 479–480 rheology, 253–254 swelling properties, 483–486 temperature dependence, dynamic storage moduli, 481–482 temperature-dependent dynamic storage moduli, 481–482 tensile properties, 480–481 Ethylene-styrene-butadiene ruber (E-SBR): pristine clay modification, rubber-clay nanocomposites, 190–191 rubber-pristine clay composites, 195 Ethylthiourea (ETU), rubber-clay nanocomposite vulcanization and permeability, 356–358 Exfoliated rubber-clay nanocomposites: apolar diene rubber, melt blending, 377

INDEX

brominated isobutylene-co-para-methylstyrene-montmorillonite nanocomposite thermodynamics, 360–362 butyl rubbers, 434–435 classification, 147–148 clay aggregation and dispersion, 454–70 formation mechanisms, 228 permeability, 347 polymer-clay morphology, 185 polymer matrices, thermodynamics model, 115–116 Fatigue behavior, reinforced elastomeric materials, 319 FENE-P model, rubber-clay nanocomposite rheology, 269–270 Filler materials: apolar diene rubbers, 370 mechanical properties, 389–392 butyl rubber clay nanocomposites: talc filler aspect ratio, 436–437 vulcanization, 454–456 clay minerals as, 35–37 reinforced elastomeric materials: clay filler mechanisms, 328–333 dynamic-mechanical (thermal) analysis, 309–310 rubber-clay nanocomposite rheology, 268–269 water-based nanocomposites, orientation, 559–561 Fire resistance: apolar diene rubber nanocomposites, 396–397 nitrile rubber composites, 424–425 Flammability, thermoplastic polyurethane nanocomposites, 495–496, 507–511 cone calorimeter test, 508–509 IR radiant panel test, 507–508 UL 94 test, 509–511 Flammability calibration, thermoplastic polyurethane nanocomposites, 495–496 Flaps (automotives), 530 Flexibility control, water-based nanocomposites, 563–566 Flocculant cations, latex compounding, rubber-clay nanocomposites, 149–153

583

Flocculants: apolar diene rubber nanocomposites, dispersion characterization, 386–387 nanocomposite morphology, 186 natural rubber latex compounding, 372–373 permeability, 347 Flory-Huggins polymer interaction, 287–288 Flory-Rehner equation, vulcanization reaction, organo-clay kinetics, 287–288 Fluoroelastomer-clay nanocomposites, rheology, 254–257 Fluorohectorite: apolar diene rubber nanocomposites: dispersion detection, 380–383 melt blending, 375–377 barrier mechanisms, 424 demand loading conditions, 420–421 natural rubber latex compounding, 372–373 olefinic rubber composites, 466–467 poly(ethylene-co-vinylacetate) (EVA) rubber-clay nanocomposite rheology, 258–259 rubber-pristine clay composites, 195, 202–206 Force field models, clay mineral mechanics, 121–123 Fracture behavior, reinforced elastomeric materials, rubber-clay nanocomposites, 322–328 cyclic loading tests, 326–328 fracture mechanics principles, 319–321 instrumented notched tensile-impact tests, 323–326 quasistatic loading, 321–323 Free energy change, melt compounding, rubber-clay nanocomposites, 158–160 Free falling dart test, reinforced elastomeric materials, 314–315 Free-radical mechanisms, sulfur vulcanization, 280–282 Freeze-drying: latex compounding, rubber-clay nanocomposites, 153 organophilic clays, 71–72

584

INDEX

Friction mechanisms: apolar diene rubber nanocomposites, 392–393 nitrile rubber composites, 421–422 Fuel hoses (automotive), rubber applications, 527 Fuel injection hosts, rubber applications, 527–528 Fuller surface optimization, reinforced elastomeric materials, 328–333 Gas permeability: automotive rubber-clay nanocomposites, 538–539 butyl rubber elastomers, 344–345 reinforcement properties, 452–454 solution method, 446–447 elastomer nanocomposites, 348–349 ethylene-propylene-diene-monomer rubber composites, 473–474 Gauche conformations: alkylammonium chains, packing density and self-assembly, 107–110 organophilic clays, cation exchange, 57–58 Geminal-dimethyl group, permeability, 344–345 Geologic properties, bentonite, 89 Gibbs free energy, brominated isobutyleneco-para-methylstyrene-montmorillonite nanocomposite thermodynamics, 360–362 Glass transition temperature, reinforced elastomeric materials, dynamicmechanical (thermal) analysis, 309–310 Glycidyl methacrylate, olefinic nanocomposites, 467–469 Grafted organoclays: direct synthesis, 62–64 olefinic nanocomposites, 467–469 solution-based, 59–61 Gum elastomers, rubber vulcanization, 276 Guth equation: apolar diene rubber nanocomposites, 391–392 nitrile rubber friction, 421 reinforced elastomeric materials, 334–335

Haloalkanes, thermal decomposition mechanism, 136 Halogenated butyl rubber, permeability, 344–345 Halpin-Tsai equations: apolar diene rubber nanocomposites, 391–392 reinforced elastomeric materials, 335 Hardness testing, reinforced elastomeric materials, 315–316 Heater hoses (automotive), rubber applications, 528 Heat flow: natural rubber vulcanization, 295–297 thermoplastic polyurethane nanocomposite flammability, 508 Heating rate tests, thermoplastic polyurethane nanocomposites, 504–507 Heat release rate (HRR): apolar diene rubber nanocomposites, fire resistance, 396–397 ethylene-propylene-diene-monomer rubber composites, 475 organo-pillared clay nanomaterials, 67–69 thermoplastic elastomers, 491 Herman’s orientation parameters, elastomer nanocomposite permeabilitgy, 351–352 Hexadecyl trimethylammonium bromide, dispersion, 386–387 High-speed tearing energy (HSTE), reinforced elastomeric materials, 324–326 Hofmann elimination, organoclay thermal decomposition, 136–137 Hofmann-Klemen (HK) effect, clay cationic exchange capacity, 27 Hoses: automotive, rubber applications, 527–528 water-based nanocomposites, 570 “House-of-cards” agglomerates: clay aggregation and dispersion, 6 latex compounding, rubber-clay nanocomposites, 152–153 Hydrogenated nitrile butadiene rubber (HNBR): ammonium cations, 209, 213–215 diene rubber composites, 219

INDEX

automotive nanocomposites, aging and ozone resistance, 539–541 barrier mechanisms, 423–424 basic characteristics, 409–410 cure characteristics, 415–416 demand loading conditions, 420–421 fire resistance, 424–425 friction/wear mechanisms, 421–422 mechanical properties, 418–421 melt compounding, 413–414 chemical factors, 163–166 compatibilizers or coupling agents, 161–163 physical factors, 166–169 pristine clay modification, primary alkenylamines, 202 vulcanization reaction, organo-clay kinetics, 285–289 Hydrophilic/lipophilic balance (HLB), organophilic clay minerals, 45–47 Hydrotalcite (HT), 4 anionic properties, 21–22 butyl rubber-clay nanocomposites, 435–436 Hydroxyl-terminated polybutadiene (HTPB) oligomer, rheology, 248–250 Illite, TOT criteria, 14 Imogolite group, TO criteria, 16 Impact loading, reinforced elastomeric materials, toughness behavior, 313–315 ImpermÒ organoclay, 65 Incompatibility, reinforced elastomeric materials and, 306–307 Inflation pressure retention (IPR), elastomer nanocomposite permeability, 346 Innerliner compound: air oven aging, 355–356 butyl rubber clay nanocomposites, 449–452 permeability measurement, 345–346, 353–356 processability, 457 In situ intercalation preparation, ethylenepropylene-diene-monomer rubber, 469–473 In situ polymerization: automotive rubber-clay nanocomposites, 536–538

585

butyl rubber clay nanocomposites, 448 elastomer nanocomposite preparation, 352 rubber-clay nanocomposites, 170 Instrumented notched tensile-impact (INIT) test, reinforced elastomeric materials, fracture mechanics, 323–326 Intercalary cations, latex compounding, rubber-clay nanocomposites, 148–149 Intercalated rubber-clay nanocomposites: apolar diene rubbers, 370–371 butyl rubbers, 434–435 classification, 147–148 clay interlayer space, 27–30 clay-polymer nanocomposites, 73–75 ethylene-propylene-diene-monomer rubber, in situ intercalation preparation, 469–473 formation mechanisms, 228–229 melt compounding, microstructure evolution, 162–163 morphology: clay modification, 184 flocculated nanocomposites, 186 polymer-clay composites, 185–186 nitrile rubber, mechanical properties, 419–421 organophilic clays, cation exchange, 49–58 crystallographic molecular structures, 54–58 solid-state intercalation, 58–59 permeability, 347 polymer chain interlayers, 229–230 rubber vulcanization, 289 solution and in situ polymerization, 170 Interface area function (IAF), reinforced elastomeric materials, 335–336 Interface-enhanced styrene-butadienerubber-clay nanocomposites, mechanical properties, 156–157 Interfacial tension: organophilic clay minerals, 46 water-based nanocomposites, 563–566 Interlayer intercalation and confinement: alkylammonium chains: packing density and self-assembly, 102–110 reversible melting transitions, 111–113

586

INDEX

Interlayer (Continued ) clay minerals, 27–30 organoclays, silane grafting, 61 rubber-clay nanocomposite formation: low molecular mass substances, 230–232 polymer chains, 229–230 Interparticle associations, clay rheology, 32–33 Interstratified clay minerals, criteria, 14–15 Intracarcass pressure (ICP), tire permeability, 354–356 Ionic exchange: clay cationic exchange capacity, 26–27 organophilic clays, cation exchange mechanisms, 53–58 Ionomers, butyl rubber nanocomposites, 439–443 IR radiant panel test, thermoplastic polyurethane nanocomposites, 507–508 Isoconversion method, thermoplastic polyurethane nanocomposites, 500–501 Isodimensional nanofillers, automotive applications, 531 Isoelectric point (IEP), clay rheology, 32–33 Isoprene-isobutadiene rubber (IIR)/clay nanocompound: apolar diene rubber nanocomposites, 392 barrier mechanisms, 396 copolymerization, 369–371 dispersion, 387–388 melt compounding, 158–160, 166–169, 377 pristine clay modification, rubber-clay nanocomposites, 190–191 solution-based preparation, 374 vulcanization, 455–456 Isoprene rubber (IR): ammonium cations, 213 apolar diene rubber nanocomposites, melt blending, 375–377 copolymerization, 369–371 permeability, 343–345 rheology, 243–246 tear behavior, 317 Kaolinite: criteria, 11 intercalation, 28–29

Kevlar-filled ethylene-propylene-dienemonomer, kinetics, 513–514 Kinetics, thermoplastic polyurethane nanocomposites, 513–516 Kissinger equation, natural rubber vulcanization, 295–297 Kraus plot, rubber-clay nanocomposite rheology, 246 Krieger’s empirical model equation, rheology of rubber-clay nanocomposites, 244–246 Labyrinth effect, apolar diene rubber nanocomposites, barrier mechanisms, 394–396 Lamellae structure, rubber-clay nanocomposite formation, 228–232 Laponite: dispersion, 8 plasma-treated clays, 69 Laser flash diffusivity, thermoplastic polyurethane nanocomposites, 497 Latex compounding method (LCM): apolar diene rubber composites, 371–373 butyl rubber nanocomposites, 447–448 future trends, 171 natural rubber-clay nanocomposite rheology, 245–246 nitrile rubber-clay composites, 411–412 rubber-clay nanocomposites, 148–157 coagulation methods, 149–153 interaction/affinity between rubber and clay, 153–157 mechanisms, 148–149 rubber-pristine clay composites, 192–194 prevulcanized latex, 194 Layered clays, rubber-clay nanocomposite morphology, multiscale organization, 184 Layered double hydroxides (LDH): automotive rubber-clay nanocomposites, 537–538 butyl rubber-clay nanocomposites, 436 clay minerals, 4 hydrotalcite, 21–22 olefinic rubber composites, 466–467 reinforced elastomeric materials, 332–333 surface area and porosity, 23–24

INDEX

Layered silicates, x-ray diffraction analysis, 182 Layer separation and miscibility, polymers, 115–121 cleavage energy, 116–121 matrix exfoliation, thermodynamics model, 115–116 surface energy, 121 Length-to-width ratio, permeability, 348 Leonov viscoelastic model, rubber-clay nanocomposite rheology, 269 Linear polymer degradation, thermoplastic polyurethane nanocomposites, 498 Lipophilicity, organophilic clay minerals, 45–47 Load-deformation diagram, reinforced elastomeric materials, 314–315 Local osmotic transition (LOT), clay swelling, 30–31 Long-chain alkenyl groups: ammonium cation substituents, 215 rubber-clay nanocomposites, ammonium cation modification, 219–228 Low molecular weight compounds: butyl rubber polymers, 444 intercalation of, 184 rubber-clay nanocomposite formation, 230–232 LPG vapor hose, rubber applications, 528 Macroscopic clay structure, 5 Maleic anhydride-grafted polymers, 443–444 adhesion properties, 456–457 ethylene-propylene-diene-monomer rubber, 479 reinforced elastomeric materials, 329–333 Mass loss rate, ethylene-propylene-dienemonomer rubber composites, 475 Masterbatch processing, organoclays: epoxidized natural rubber, 414 thermal decomposition mechanism, 136–137 Maximum torque value, rubber vulcanization, curing mechanics, 278 Maxwell-Wagner polarization effects, nitrile rubber composites, 425

587

MDR rheomtery, innerliner compound vulcanization, 356 Mechanical properties: apolar diene rubber nanocomposites, 387–392 clays and clay minerals, 121–123 ethylene-propylene-diene-monomer rubber, 476–479 nitrile rubber composites, 416–421 reinforced elastomeric materials, 307–322 compression set determination, 317–319 dynamic-mechanical analysis, 307–310 fatigue behavior, 319 fracture mechanics, 319–321 hardness testing, 315–316 tear behavior, 316–317 tensile testing, 310–313 toughness behavior, impact loading conditions, 313–315 rubber vulcanization, 290–297 Melt compounding method (MCM), rubberclay nanocomposites, 55–170 apolar diene rubber nanocomposites, 374–377 butyl rubbers, 445 future trends, 171 hydrogenated nitrile butadiene rubber, pristine clay modification, 202 mechanisms, 157–160 mixing curative effects, 162 nitrile rubber-clay composites, 412–414 organic modification, matrix rubber vulcanization, 169–170 permeability, 352 pristine clay distribution and dispersion, 190–191 aggregated composites, 194–197 coagulating agents, 196–197 rubber-pristine clay composites, 192–194 thermoplastic clay nanocomposites comparisons, 160–162 vulcanization and microstructure evolution, 162–169 Mercaptobenzothiazole (MBT), rubber-clay nanocomposite vulcanization and permeability, 356–358

588

INDEX

Mesoscale, clays and clay materials, 5 Mesoscopic rheological model, rubber-clay nanocomposites, 269–270 Methylene blue (MB) adsorption, clay surface area and porosity, 23–24 Methyl ethyl ketone (MEK), nitrile rubberclay composites, 410–411 Micas: alkylammonium chains: packing density and self assembly, 104–110 reversible melting transitions, 111–113 butyl rubber clay nanocomposites, 437 TOT criteria, 14 Micromechanical models, reinforced elastomeric materials, 334–335 Microstructure evolution: melt compounding, rubber-clay nanocomposites, rubber vulcanization and, 162–163 thermoplastic polyurethane nanocomposites, 511–512 Microwave treatment, organophilic clays, 71–72 Minerals, clays as, 3–4 Minimum torque value, rubber vulcanization, curing mechanics, 278 Mining, bentonite, 89–90 Modulus reduction factor (MRF): nitrile rubber composites, 421 reinforced elastomeric materials, 335 Molecular dynamics, thermoplastic polyurethane nanocomposites, 497–498 Molecular scale, clays and clay materials, 5 Molecular simulation, alkylammonium chains, 111 Monolayers, organophilic clays, cation exchange, 54–58 Monte Carlo simulation, rheology of rubberclay nanocomposites, 245–246 Montmorillonite: alkylammonium chains: packing density and self assembly, 102–110 reversible melting transitions, 112–114 apolar diene rubber nanocomposites: melt blending, 374–377 solution-based preparation, 373–374

BIMSM nanocomposites, permeability, 351–352 brominated isobutylene-co-para-methylstyrene nanocomposites, thermodynamics, 358–362 butyl rubber-clay nanocomposites, 435 cleavage energy, 116–121 commercial products, 65–66 ethylene-propylene-diene-monomer rubber, in situ intercalation preparation, 469–473 fluoroelastomer-clay nanocomposites, rheology, 254–257 layered structure, 10–11 melt compounding, rubber-clay nanocomposites, thermoplastic-clay nanocomposites and, 160–162 olefinic rubber composites, 466–467 reinforced elastomeric materials, 330–333 tensile testing, 312–313 reinforced elastomeric materials and, 306–307 rubber-clay nanocomposites: ammonium cations, 207–212 long-chain alkenyl substituents, 219–228 polar group ammonium cations, 215–219 pristine clay modifiers, 182–183, 190–191 substituent modification, 212–215 vulcanization activation, 139–141 rubber-pristine clay composites, 193–194 melt blending, 194–195 primary alkenylamines, 198–202 solution-based, 196 solid-state intercalation, 59 styrene-butadiene-rubber-clay nanocomposites, rheology, 246–247 surface tension and cleavage energies, 116–122 swelling properties, 17, 19–20 thermoplastic polyurethane nanocomposites, 493 vulcanization reaction, organo-clay kinetics, 286–289 Mooney-Rivlin plot, apolar diene rubber nanocomposites, 389–392

INDEX

Mooney viscosity, brominated isobutyleneco-para-methylstyrene-montmorillonite nanocomposite thermodynamics, 358–362 Morphology, rubber-clay nanocomposites: ammonium cation modification, 206–212 basic principles, 181 butyl rubbers, 433–435 cationic clay properties, 182–183 clay distribution and dispersion, 183 long-chain alkenyl substituents, ammonium cation-modified montmorillonite, 219–228 low molecular mass substances, 183 montmorillonite modification, 212–215 multiscale organization, layered clays, 183 polar group ammonium cations, montmorillonite modification, 215–219 polymer-clay composites, 183–186 primary alkenylamines, 197–206 pristine clays, 186–197 proposed mechanisms, 228–232 thermoplastic polyurethanes, 494–495 x-ray diffraction analysis, 182 Multiaxial deformation behavior, reinforced elastomeric materials, 314–315 Multifunctional additives (MFA), rubberclay nanocomposites, vulcanization activation, 138–141 Multiscale organization: clays and clay minerals, 6–8 delamination/exfoliation vs. stacking, 6–8 dispersion vs. aggregation, 6 rubber-clay nanocomposite morphology, layered clays, 184 m-xylylenediamine (MXD), latex compounding, rubber-clay nanocomposites, 155–157 Nanocor organoclay, 65 Nanofil compounds, cyclic loading, fracture mechanics testing, 327–328 Nanomaterials: clay-polymer morphology and classification, 185–186 clays as, 20–21 mechanical data, 121, 123

589

organo-pillared clays, 66–69 plasma-treated clays, 69 organoclay nanocomposites, 72–75 nanoMaxÔ, 65 Nanopolymers, clays as, 20–21 Nanoscale, clays and clay materials, 5 Natural butadiene rubber (NBR): automotive rubber-clay nanocomposites, 535–538 resistance, 541–542 barrier mechanisms, 424 basic characteristics, 409–410 diene rubber composites, 225 fire resistance, 424–425 friction/wear mechanisms, 421–422 latex compounding, 411–412 mechanical properties, 416–421 nitrile rubber-clay composites: chloroform solvents, 411 cure characteristics, 414–416 vulcanization, 287–288 Natural clays, availability, 33–35 Natural rubber (NR): ammonium cations, 213 apolar diene rubber nanocomposites, 393 barrier mechanisms, 394–396 automotive rubber-clay nanocomposites, 535–538 diene rubber composites, 221–222 latex compounding, 371–373 mechanical properties, 387–392 melt blending process, 374–377 polar group ammonium cation modification, diene rubber composites, 217 pristine clay modification, rubber-clay nanocomposites, 190–191 primary alkenylamines, 198–200 reinforced elastomeric materials: cyclic loading, fracture mechanics testing, 327–328 fracture mechanics, 325–326 rubber-clay nanocomposites: ammonium cations, 207 rheology, 243–246 solution-based preparation, 373–374 sulfur vulcanization, 278–282 tear behavior, 317 vulcanization kinetics, 292–297

590

INDEX

Nanomaterials (Continued ) water-based nanocomposites, tennis balls, 567–568 Natural rubber/chloroprene (NR/CR) blends, vulcanization reaction, organo-clay kinetics, 284–289 Natural rubber-ethylene natural rubber blends, rubber-pristine clay composites, 195 Negatively charged clay layers, 4 Neutral clays, layer organization, 8–11 Newtonian behavior, rheology of rubber-clay nanocomposites, 244–246 Nitrile rubber rubber-clay nanocomposites: barrier properties, 423–424 basic principles, 409–410 cure characteristics, 414–416 fire resistance, 424–425 friction/wear properties, 421–422 future research issues, 425–426 mechanical properties, 416–421 preparation methods and dispersion, 410–414 latex compounding, 411–412 melt blending, 412–414 solution-base methods, 410–411 Nitrogen physisorption: clay aggregation and dispersion, 8 nitrile rubber composites, 423–424 Nonphyllosilicate cationic layered minerals, layered structure, 21 Nucleophilic substitution, rubber-clay nanocomposites, vulcanization activation, 139–141 Octadecylamine (ODA): apolar diene cure characteristics, 377–379 latex compounding, rubber-clay, 157 melt compounding, rubber-clay nanocomposites, vulcanization effects, 169–170 nitrile rubber-clay composites, melt mixing, 413–414 rubber-pristine clay composites, primary alkenylamines, 197–206 Octadecylammonium (ODA), poly(ethyleneco-vinylacetate) (EVA) rubberclay nanocomposite rheology, 258–259

Octadecyltrimethylammonium mica, reversible melting transitions, 111–113 Octahedral sheets: anionic clay minerals, 21–22 clay layer organization, 9–11 nonphyllosilicate cationic layers, 21 Olefin rubber-clay nanocomposites: basic properties, 465–466 compatibilizer systems, 467–469 creep properties, 482–483 gas permeability, 473–474 mechanical properties, 476–479 mineral compounds, 466–467 organoclay thermal decomposition, 136–137 preparation, 479–480 rheological properties, 474 in situ intercalation method, 469–473 stability, 475 swelling properties, 475–476, 483–486 temperature dependence, dynamic storage moduli, 481–482 tensile properties, 480–481 One-pot clay-polymer nanocomposite formation, 66 Onsager orientation distribution, elastomer nanocomposite permeabilitgy, 351–352 Onset decomposition temperature, organoclay thermal decomposition, 133–135 Optimal clay concentration, clay mineral purification, 94 Optimum cure time, rubber vulcanization, curing mechanics, 278 Orbital ball-on-plate test, nitrile rubber composites, 421–422 Organically modified montmorillonite (OMMT): butyl rubber nanocomposites, solution method, 445–447 reinforced elastomeric materials, dynamicmechanical (thermal) analysis, 309–310 Organic cation decompostion, rubber-clay nanocomposites, 128–135 Organic clay minerals: automotive applications, 534 chemical treatments, 71

INDEX

clay-polymer nanocomposites, 72–75 commercial products, 64–66 dispersion maximization, 69–70 hydrophilic/lipophilic balance, 45–47 melt compounding, rubber-clay nanocomposites, 158–160 thermoplastic-clay nanocomposite comparisons, 160–162 vulcanization effects, 169–170 nanomaterials, 66–69 one-pot CPN formation, 66 physical treatments, 71–72 polymer technology, 47–49 postsynthesis modifications, 64 rubber-clay nanocomposites: cation decompostion, 128–135 vulcanization activation, 137–141 solid-state intercalation, 58–59 solution-based grafting, 59–61 synthesis procedures, 49–66 thermal decomposition mechanism, 135–137 thermal stability, 70–71 vulcanization reaction, 283–289 Organic-inorganic interface, alkylammonium chains, packing density and self-assembly, 108–110 Organic montmorillonite (OMMT) nanocomposites: apolar diene rubber nanocomposites, 390–392 butyl rubber nanocomposites, 439 reinforced elastomeric materials, 330–333 compression set determination, 318–319 solution and in situ polymerization, 170 Organic species, clay intercalation, 28 Organic substances, clay reactions, 97–99 Organo-pillared clay nanomaterials, 66–69 Orientation parameters, water-based nanocomposites, barrier mechanisms, 558–561 Oscillating disc rheometer (ODR), rubber vulcanization, curing mechanics, 277–278, 290–297 Osmotic transition (OT), clay swelling, 30–31

591

Ozawa equation, natural rubber vulcanization, 295–297 Ozone cracking time, apolar diene rubber nanocomposites, 397–398 Ozone resistance, automotive rubber-clay nanocomposites, 539–541 Packing density, alkylammonium chains, 102–110 Pade approximation, rheology of rubber-clay nanocomposites, 245–246 Palygorskite, TOT criteria, 15–16 Paraffin-like structures, organophilic clays, cation exchange, 54–58 Particle separation: latex compounding, rubber-clay nanocomposites, 148–153 melt compounding, rubber-clay nanocomposites, 163 Particle size: clays and clay materials, 5 organic clays, modification, 99 reinforced elastomeric materials, 331–333 Particulate dispersion, elastomer nanocomposite permeability, 346 Payne effect: reinforced elastomeric materials, dynamicmechanical (thermal) analysis, 309–310 rubber-clay nanocomposites, rheological behavior, 242, 263–268 PellethaneÔ, 491–493 dispersion, 501–503 kinetics, 514–516 thermal properties, 504–507 Percolation threshold, apolar diene rubber nanocomposites, 391–392 Performance evaluation, rubber-clay nanocomposite permeability, 362–364 Permeability: automotive rubber-clay nanocomposites, 538–539 butyl rubber clay nanocomposites, 449–452 rubber-clay nanocomposites: basic principles, 343 BIMSM montmorillonite thermodynamics, 358–362

592

INDEX

Permeability (Continued ) butyl rubber elastomers, 343–345 elastomer nanocomposites, 352 exfoliated nanocomposites, 347 flocculated nanocomposites, 347 gas permeability, 348 intercalated-flocculated nanocomposites, 347 intercalated nanocomposites, 347 Nielsen’s model, 348 Onsager orientation distribution function, 351 order parameters, 348–349 orientation parameter, 349–350 plate formation, 347–352 small-angle x-ray scattering, 349–350 temperature and compound materials, 352–356 tire innerliner compound, 345–346 tire permeability, 346, 362–364 tortuosity effect, 347–348 vulcanization effects, 356–358 Permeation behavior, reinforced elastomeric materials, compression set determination, 318–319 Peroxide vulcanization, curing kinetics, 282–283 Phenolic resins, rubber-pristine clay composites, 192 Phenomenological kinetics, rubber vulcanization, 290–297 Phosphine ionomer, butyl rubber nanocomposites, 439–443 Phyllosilicates, TOT criteria, 14 Physical treatments, organophilic clays, 71–72 Physisorption isotherm, clay surface area and porosity, 22–24 Pin-on-plate (POP) sliding wear test, nitrile rubber friction/wear mechanisms, 421–422 Plasma-treated clays, nanomaterials, 69 Platelet aggregation: clays and clay minerals, 6 rubber-clay nanocomposites, basic properties, 241–242 Plate orientation, elastomer nanocomposite gas permeability, 349–351

Pneumatic devices, water-based nanocomposites, 566–570 basketballs, 569 bicycle and specialty tires and tubes, 570 soccer balls, 568–569 tennis balls, 567–568 Point of zero charge (PZC), clay rheology, 32–33 Polar group ammonium cations, montmorrillonite rubber composites, 215–219 Polyepichlorohydrin rubber-clay nanocomposites, rheology, 259–261 Poly(ethylene-co-vinylacetate) (EVA) rubber-clay nanocomposite, pristine clay modification, 257–259 Polyethylene terephthalate (PET), reinforced elastomeric materials, 330–333 Polymer nanocomposites: automotive applications, 531 intercalation, 29–30 layer separation and miscibility, 115–121 cleavage energy, 116–121 matrix exfoliation, thermodynamics model, 115–116 surface energy, 121 low molecular weight butyl rubbers, 444 organophilic clays, 47–49 reinforced elastomeric materials, chain behavior, 329–333 resins, thermoplastic elastomer nanocomposites, 491–493 rubber-clay nanocomposites: chain intercalation, 229–230 morphology and classification, 184–186 organic cation decompostion, 128–135 Polysulfide ions, rubber-clay nanocomposites, vulcanization activation, 140–141 Polyurethane: automotive nanocomposites, 544–545 automotive rubber-clay nanocomposites, 535–538 natural rubber latex compounding, 373 thermoplastic-polyurethane-clay nanocomposites, 261–262 Porosity, clay minerals, 22–24 Porous clay heterostructures (PCH), postsynthesis modifications, 64 Positively charge clay layers, 4

INDEX

Postsynthesis modification, grafted organoclays, 62–64 Power-law index, fluoroelastomer-clay nanocomposite rheology, 254–257 Power steering return hose, rubber applications, 528 Precursor organic cation salts, organoclay thermal decomposition, 133–135 Preshear parameters, isopropylene rubberclay nanocomposite rheology, 245–246 Pressure accumulators, water-based nanocomposites, 570 Preswelled organic clay materials, melt compounding, rubber-nanoclay composites, physical factors, 169 Prevulcanized latex, rubber-pristine clay composites, 194 procedures for, 197 Primary alkenylamines, rubber-pristine clay composites, 197–206 fluorohectorite modification, 200–206 montmorillonite and bentonite, 198–200 Primary amine (PRIM), olefinic nanocomposites, 468–469 Primary organic montmorillonite (P-OMt), apolar diene rubber nanocomposites, 390–392 Pristine clay, rubber-clay nanocomposite morphology: aggregated clay organization, rubber materials, 189–190 cationic clays, 187 distribution and dispersion, 187–189 latex composites, 192–194 matrix distribution and dispersion, 190–194 melt compounding, 190–191 organic modifiers, 182–183 rubber solution composites, 192 Processing methods, rubber-clay nanocomposites: automotive elastomers, 542–544 butyl rubbers, 457 future trends, 170–171 latex compounding, 148–157 coagulation methods, 149–153 interaction/affinity between rubber and clay, 153–157

593

mechanisms, 148–149 melt compounding, 157–170 mechanisms, 157–160 mixing curative effects, 162 organic modification, matrix rubber vulcanization, 169–170 thermoplastic clay nanocomposites comparisons, 160–162 vulcanization and microstructure evolution, 162–169 overview, 147–148 solution intercalation and in situ polymerization intercalation, 170 thermoplastic elastomers, 491 Proton exchange membrane fuel cells, plasma-treated clays, 69 Pseudoplastic behavior, automotive elastomers, 542–544 Pseudotrimolecular arrangement, organophilic clays, cation exchange, 57–58 Purification, clays and clay minerals, 93–97 clay concentration, 94 swelling time, 94–95 temperature, 95–97 Pyrophyllite, TOT criteria, 14 Quasistatic loading, reinforced elastomeric materials, fracture mechanics, 321–323 Quaternary amine (QUAT), olefinic nanocomposites, 468–469 Quaternary ammonium cations, thermal decomposition mechanism, 135–137 Quaternary organic montmorillonite (P-OMt), apolar diene rubber nanocomposites, 390–392 Quaternary phosphonium salts, organophilic clays, rubber vulcanization, 289 Radiant panel testing, thermoplastic polyurethane nanocomposites, 496 microstructure evolution, 511–512 R-curves, reinforced elastomeric materials, fracture mechanics, 322–323 Reactive chamber pattern, apolar diene rubber nanocomposites, melt blending, 375–377

594

INDEX

Reinforced elastomeric materials, rubberclay nanocomposites: basic principles, 305–307 butyl rubber nanocomposites, 452–454 fracture behavior, 322–328 cyclic loading tests, 326–328 fracture mechanics principles, 319–321 instrumented notched tensile-impact tests, 323–326 quasistatic loading, 321–323 mechanisms, 328–333 theories and modeling, 333–334 viscoelastic and mechanical testing, 307–322 compression set determination, 317–319 dynamic-mechanical analysis, 307–310 fatigue behavior, 319 hardness testing, 315–316 tear behavior, 316–317 tensile testing, 310–313 toughness behavior, impact loading conditions, 313–315 Relaxation time, polyepichlorohydrin rubber-clay nanocomposite rheology, 259–261 Resins: low molecular mass substances, 444 thermoplastic elastomer nanocomposites, 491–493 Rheology: clay materials, 31–33 ethylene-propylene-diene-monomer composites, 474 rubber-clay nanocomposites: acrylonitrile butadiene rubber, 250–253 basic principles, 241–242 data thresholds and experiments, 263–269 epoxidized natural rubber and polyisoprene rubber, 243–246 ethylene-propylene nanocomposites, 253–254 fluoroelastomer-clay, 254–257 future research, 270 polybutadiene rubber, 247–250 polyepichlorohydrin, 259–261

poly(ethylene-co-vinylacetate) (EVA), 257–259 poly(isobutylene-co-para-methylstyrene) (BIMS), 257 polymer-clay nanocomposites, 269–270 styrene-butadiene composites, 246–247 styrene-ethylene-buylene-styrene (SEBS) block copolymer-clay nanocomposites, 262–263 thermoplastic polyurethane, 261–262 Rolling-on-plate (ROP) configuration, nitrile rubber composites, 421–422 Rubber boots (automotive), 529 Rubber-clay nanocomposites: automotive applications: aging and ozone resistance, 539–541 barrier/air retention properties, 538–539 basic principles,525–526 disadvantages, 548–549 elastomeric components, 531 elastomeric nanocomposites, 531–548 organoclay minerals, 534 polyurethane-organoclay composites, 544–545 processability, 542–544 rubber, 526–531 rubber-clay vs. other nanofillers, 534 solvent resistance, 541–542 tire nanocomposites, 545–548 weight and balanced mechanical properties, 535–538 vulcanization reaction, organo-clay kinetics, 283–289 Rubber-clay nanocomposites (RCNs): bentonite, 87–89 chemistry: basic principles, 127–128 microstructure evolution, 163–169 organic cation decomposition, 128–135 rubber vulcanization activation, organic cations, 137–141 thermal decomposition mechanism, 135–137 morphology: ammonium cation modification, 206–212

INDEX

basic principles, 181 cationic clay properties, 182–183 clay distribution and dispersion, 183 long-chain alkenyl substituents, ammonium cation-modified montmorillonite, 219–228 low molecular mass substances, 183 montmorillonite modification, 212–215 multiscale organization, layered clays, 183 polar group ammonium cations, montmorillonite modification, 215–219 polymer-clay composites, 183–186 primary alkenylamines, 197–206 pristine clays, 186–197 proposed mechanisms, 228–232 x-ray diffraction analysis, 182 permeability: basic principles, 343 BIMSM montmorillonite thermodynamics, 358–362 butyl rubber elastomers, 343–345 elastomer nanocomposites, 352 exfoliated nanocomposites, 347 flocculated nanocomposites, 347 gas permeability, 348 intercalated-flocculated nanocomposites, 347 intercalated nanocomposites, 347 Nielsen’s model, 348 Onsager orientation distribution function, 351 order parameters, 348–349 orientation parameter, 349–350 plate formation, 347–352 small-angle x-ray scattering, 349–350 temperature and compound materials, 352–356 tire innerliner compound, 345–346 tire permeability, 346, 362–364 tortuosity effect, 347–348 vulcanization effects, 356–358 physical properties, chemistry, 163–169 processing methods: future trends, 170–171 latex compounding, 148–157 melt compounding, 157–170 overview, 147–148

595

solution intercalation and in situ polymerization intercalation, 170 reinforcing components: basic principles, 305–307 elastomer fracture behavior, 322–328 mechanisms, 328–333 theories and modeling, 333–334 viscoelastic and mechanical properties, 307–322 rheology: acrylonitrile butadiene rubber, 250–253 basic principles, 241–242 data thresholds and experiments, 263–269 epoxidized natural rubber and polyisoprene rubber, 243–246 ethylene-propylene nanocomposites, 253–254 fluoroelastomer-clay, 254–257 future research, 270 polybutadiene rubber, 247–250 polyepichlorohydrin, 259–261 poly(ethylene-co-vinylacetate) (EVA), 257–259 poly(isobutylene-co-para-methylstyrene) (BIMS), 257 polymer-clay nanocomposites, 269–270 styrene-butadiene composities, 246–247 styrene-ethylene-buylene-styrene (SEBS) block copolymer-clay nanocomposites, 262–263 thermoplastic polyurethanes, 261–262 vulcanization: basic principles, 275–276 kinetics, 290–297 organoclay activation, 137–141, 283–289 reaction mechanics, 276–278 rubber cross-linking systems, 278–283 Rubber matrices, pristine clays: aggregation, 189–190 distribution and dispersion, 187–189 Rubber solution, rubber-clay nanocomposites, 192 Salts, rubber-clay nanocomposites, organic cation decompostion, 128–135

596

INDEX

Scanning electron microscopy (SEM), apolar diene rubber nanocomposites, dispersion detection, 380–383 Schallamach pattern, nitrile rubber composites, 421–422 “Schizophrenic molecules,” organophilic clay amphophilicity, 46 Scorch time: rubber vulcanization, curing mechanics, 278 sulfur vulcanization, 279–282 Seals (automotive), rubber applications, 528–529 Self-assembly, alkylammonium chains, 102–110 Semi-rigid clay networks, reinforced elastomeric materials, 333 Separated phase microcomposites, morphology and classification, 185 Separated rubber-clay nanocomposites, 147–148 Sepiolite: butyl rubber clay nanocomposites, 438 TOT criteria, 15 Serpentine, criteria, 11 Shear viscosity: butadiene-rubber nanocomposite rheology, 249–250 polyepichlorohydrin rubber-clay nanocomposites, 259–261 rubber-clay nanocomposite rheology, 268–269 Shore A hardness: reinforced elastomeric materials, 315–317 thermoplastic polyurethane nanocomposites, 504–507 Short-chain alkenyl groups, ammonium cation substituents, 215 Silane grafting, organoclays, 59–61 Silicates: apolar diene rubbers, 370–371 elastomer nanocomposite gas permeability, 348–349 nitrile rubber-clay composites, 410 pristine clay modification, rubber-clay nanocomposites, 191 rubber-pristine clay composites, 195 rubber vulcanization, 290–297

Single-edge-notched tension (SENT) specimens, reinforced elastomeric materials, fracture mechanics, 321–323 Size parameters, latex compounding, rubberclay nanocomposites, 149–157 Slurry additives, rubber-pristine clay composites, 193 coagulating agents, 196–197 Small angle X-ray scattering (SAXS): apolar diene rubber nanocomposites, dispersion characterization, 386–387 elastomer nanocomposite gas permeability, 349–350 Small molecules, clay intercalation, 27–28 Smectites: rheology, 31–33 swelling properties, 17, 19–20 TOT criteria, 14 SN2 substitution reaction, organoclay thermal decomposition, 137 Soccer balls, water-based nanocomposites, 568–569 Sol-gel transition (SGT), clay swelling, 30–31 Solid-state intercalation, organophilic clays, cation exchange, 58–59 Solution-based cation exchange, organophilic clays, 49–58 Solution intercalation: apolar diene rubber processing, 373–374 butyl rubbers, 445–447 nitrile rubber-clay composites, 410–411 rubber-clay nanocomposites, 170, 192 permeability, 352 Solvents: apolar diene rubber processing, 373–374 automotive nanocomposites, resistance, 541–542 evaporation, alkyl surfactants, 114–115 Sonication, organophilic clays, 71–72 Spatial distribution, melt compounding, rubber-clay nanocomposites, 163 Specific heat analysis, thermoplastic polyurethane nanocomposites, 497 Specific surface area (SSA): clay aggregation and dispersion, 70 clay surface area and porosity, 22–24

INDEX

Sports balls, water-based nanocomposites, 566–570 basketballs, 569 bicycle and specialty tires and tubes, 570 soccer balls, 568–569 tennis balls, 567–568 Stability, ethylene-propylene-diene-monomer rubber, 475 Stacking, clay aggregation and dispersion, 6–8 Staining, water-based nanocomposites, 572 Star-branched butyl (SBB) polymers: automotive rubber-clay nanocomposites, 538 characteristics, 431–433 Stiffness, reinforced elastomeric materials, 335–336 Storage modulus: ethylene-propylene monomer nanocomposites, temperature dependence, 481–482 isopropylene rubber-clay nanocomposite rheology, 245–246 poly(ethylene-co-vinylacetate) (EVA) rubber-clay nanocomposite rheology, 258–259 reinforced elastomeric materials, dynamicmechanical (thermal) analysis, 308–310 rubber-clay nanocomposite rheology, 268–269 Strength testing, reinforced elastomeric materials, tensile testing, 310–313 Stress-strain analysis, styrene-butadienerubber-clay nanocomposites, 155–157 Structure-property relationships, thermoplastic elastomers, 491 Styrene-butadiene-rubber-clay nanocomposites: ammonium cations, 209, 213 diene rubber composites, 217, 219 diene rubber composites, 225 latex compounding, 155–157 melt compounding, 159–160 pristine clay modification, primary alkenylamines, 200–202 Styrene-butadiene-rubber (SBR)-clay nanocomposites:

597

aging and ozone resistance, 539–541 apolar diene rubber nanocomposites: barrier mechanisms, 395–396 dispersion, 386–387 mechanical properties, 391–392 melt blending, 376–377 automotive rubber-clay nanocomposites, 535–538 copolymerization, 369–371 reinforced elastomeric materials: fracture mechanics, 325–326 tensile testing, 311–313 toughness behavior, 313–315 rheology, 246–247 solution-based preparation, 374 Styrene-ethylene-butylene-styrene (SEBS) block copolymer-clay nanocomposites, rheology, 262–263 Substituent modification, montmorillonite rubber-clay nanocomposites, 212–215 Substitution reactions, rubber-clay nanocomposites, vulcanization activation, 140–141 Sulfur curatives: apolar diene rubbers, mechanical properties, 389–392 melt compounding, rubber-nanoclay composites, 166–169 olefinic nanocomposites, 467–469 vulcanization, 278–282 Surface active ingredients, rubber-clay nanocomposites, vulcanization activation, 138–141 Surface area, clay minerals, 22–24 Surface energy, polymer layer separation and miscibility, 121 Surface temperature, thermoplastic polyurethane nanocomposite flammability, 508 Surfactant/CEC ratio: alkylammonium chains: diffusion dynamics, 110 solvent evaporation and thermal elimination, 113–115 clay-polymer nanocomposites, 73–75 organophilic clays, dispersion maximization, 69–70

598

INDEX

Surfactant properties: butyl rubber nanocomposites, 439 latex compounding, rubber-clay nanocomposites, 155–157 Swelling clays: basic principles, 30–31 butyl rubber nanocomposites, 439 ethylene-propylene-diene-monomer rubber composites, 475–476 ethylene-propylene monomer nanocomposites, 483–486 interlayer intercalation and confinement, 27–30 layered structure, 17, 22–23 polymer intercalation, 29–30 Swelling time, clay mineral purification, 94–95 Synthetic clays, availability, 33–35 Talc fillers: butyl rubber clay nanocomposites, 436–437 TOT criteria, 14 Tallow groups, ammonium cations, 213–215 diene rubber composites, 220–228 Tear and fatigue analyzer (TFA), reinforced elastomeric materials, cyclic loading, fracture mechanics testing, 326–328 Tear behavior, reinforced elastomeric materials, 316–317 Tear strength, reinforced elastomeric materials, 331–333 Temperature: butyl rubber nanocomposites, air permeability, 452 clay mineral purification, 95–97 ethylene-propylene monomer nanocomposites, dynamic storage moduli, 481–482 nanocomposite permeability, 352–356 natural rubber vulcanization, 295–297 Tennis balls, water-based nanocomposites, 567–568 Tensile testing: apolar diene rubber nanocomposites, 389–392 ethylene-propylene-diene-monomer rubber, 476–479

reinforced elastomeric materials, 310–313 Ternary systems, clay-polymer nanocomposites, 73–75 Tetraalkylammoniums, solid-state intercalation, 59 Tetrahedral sheets: cationic layered silicates, 21 clay layer organization, 9–11 Tetramethylthiuram disulfide (TMTD), vulcanization and permeability, 356–358 Thermal properties: alkylammonium chains, 111–115 reversible melting transitions, 111–113 solvent evaporation and thermal elimination, 113–115 alkyl surfactants, 114–115 organophilic clays, 70–71 rubber-clay nanocomposites: organic cation mechanism, 128–135 organoclays, 135–137 thermoplastic polyurethane nanocomposites, 495, 503–507 diffusivity, 497 Thermodynamics, brominated isobutyleneco-para-methylstyrene-montmorillonite nanocomposite, 358–362 Thermogravimetric analysis (TGA): thermoplastic elastomers, 489–491, 495 thermoplastic polyurethane nanocomposites: decomposition kinetics, 499–501 thermal properties, 504–507 Thermophysical properties, thermoplastic polyurethane nanocomposites, 496–497, 512–513 Thermoplastic elastomer (TPE) nanocomposites: automotive applications, 544–545 basic properties, 489–491 calibration, 495 decomposition kinetics modeling, 497–501 experimental analysis, 495 flammability properties, 495–496, 507–511 kinetic parameters, 512–516 materials selection, 491–493

INDEX

melt compounding, 160–162 microstructures, 511–512 morphology, 494–495 nanoparticles, 493 dispersion, 501–503 thermal properties, 503–507 physical properties, 496–497, 512–513 polymer resins, 491–493 processing, 493–494 water-based nanocomposites vs., 566 Thermoplastic polyurethane nanocomposites (TPUN): basic properties, 489–491 flammability, 495–496, 507–511 cone calorimeter test, 508–509 IR radiant panel test, 507–508 UL 94 test, 509–511 kinetics, 513–516 microstructure analysis, 511–512 morphology, 494–495 nanoparticles, 493 resins, 491–493 rheology, 261–262 thermophysical properties, 496–497, 503–507, 512–513 Thiuram vulcanization accelerator, melt compounding, rubber-clay nanocomposites, 164–169 Thixotropy, clay rheology, 33 Tires: bicycle and specialty tires, water-based nanocomposites, 570 butyl rubber clay nanocomposites, 452 reinforcement properties, 453–454 improvement in, 346 innerliner compound, 345–346 nanocomposite performance evaluation, 362–364 organoclay nanocomposites, 545–548 rubber applications, 529–530 TO (tetrahedral-octahedral) clay groups, 11–17 nonphyllosilicate cationic layers, 21 Toluene, apolar diene rubber processing, 373–374 Torque delta: natural vulcanization, 292–297 reinforced elastomeric materials, 331–333

599

Tortuosity effect: apolar diene rubber nanocomposites, barrier mechanisms, 394–396 elastomer nanocomposite permeability, 347–348 nitrile rubber composites, 424 water-based nanocomposite modeling, 561–563 TOT (tetrahedral-octahedral-tetrahedral) clay groups, 11–17 surface area and porosity, 23–24 Toughness behavior, reinforced elastomeric materials, impact loading conditions, 313–315 Trans conformations, organophilic clays, cation exchange, 57–58 Transmission electron microscopy (TEM): apolar diene rubber nanocomposites, dispersion detection, 380–383 organophilic clay nanocomposities, 73–75 thermoplastic polyurethane nanocomposites, 489–491 Tributyl phosphate (TMt), fire resistance, 396–397 Bis-(3-Triethoxysilylpropyl)-tetrasulfide (TESPT): apolar diene rubber nanocomposites, 391–392 rheology, 246–247 Triisopropanolamine (TA), latex compounding, rubber-clay nanocomposites, 155–157 Trimethylammonium, reversible melting transitions, 111–113 Trimethyloctadecylammonium chloride (TMO), organoclay thermal decomposition, 133–135 Trioctahedral smectites, swelling properties, 17, 19–20 Tubing (automotive), rubber applications, 529–530 UL 94 test, thermoplastic polyurethane nanocomposites, 496, 509–510 posttest microstructures, 512 Unsaturated organic ammonium chloride (UOAC), latex compounding, rubberclay nanocomposites, 155–157

600

INDEX

Van der Waals interactions: alkylammonium chains, cleavage energy, 119–121 butyl rubber elastomers, 345 Vermiculites: butyl rubber clay nanocomposites, 437 latex compounding, 448 TOT criteria, 14 Viscoelastic testing, reinforced elastomeric materials, 307–322 compression set determination, 317–319 dynamic-mechanical analysis, 307–310 fatigue behavior, 319 hardness testing, 315–316 tear behavior, 316–317 tensile testing, 310–313 toughness behavior, impact loading conditions, 313–315 Viscosity: butadiene-rubber nanocomposite rheology, 249–250 fluoroelastomer-clay nanocomposite rheology, 254–257 Mooney viscosity, brominated isobutylene-co-para-methylstyrene-montmorillonite nanocomposite thermodynamics, 358–362 poly(ethylene-co-vinylacetate) (EVA) rubber-clay nanocomposite rheology, 258–259 rubber-clay nanocomposites, rheological behavior, 242 Viscosity percolation threshold, rheology of rubber-clay nanocomposites, 244–246 Vulcanization: brominated isobutylene-co-paramethylstyrene-montmorillonite nanocomposite thermodynamics, 358–362 butyl rubber clay nanocomposites, 454–456 ethylene-propylene-diene monomers, 469–474 innerliner compounds, 355–356 melt compounding, rubber-clay nanocomposites: microstructure evolution, 162–163 OMC influence on, 169–170

nitrile-rubber composites, cure characteristics, 414–416 rubber-clay nanocomposites: basic principles, 275–276 kinetics, 290–297 organoclay activation, 137–141, 283–289 reaction mechanics, 276–278 rubber cross-linking systems, 278–283 vulcanization, 356–358 Water-based nanocomposites: barrier properties, 557–566 dispersion and orientation, 558–561 interfacial interactions and flexibility control, 563–565 tortuous path modeling, 561–563 blooming and staining reduction, 572 breakthrough time applications, 571 chemical protective gear, 571–572 hoses, 570 pressure accumulators, 570 sports balls and pneumatic applications, 566–570 basketballs, 569 bicycle and specialty tires and tubes, 570 soccer balls, 568–569 tennis balls, 567–568 thermally processed elastomers vs., 566 Water removal, latex compounding, rubberclay nanocomposites, 152–153 Waves of detachment, apolar diene rubber nanocomposites, 392–393 Wear mechanisms: apolar diene rubber nanocomposites, 392–393 nitrile rubber composites, 421–422 Weight loading: automotive rubber-clay nanocomposites, 535–538 thermoplastic polyurethane nanocomposite flammability, 507–508 Weight loss data, thermoplastic polyurethane nanocomposites, thermal properties, 504–507 Wide-angle x-ray diffraction (WAXD) peaks: latex compounding, rubber-clay nanocomposites, 149–153

INDEX

melt compounding, rubber-nanoclay composites, physical factors, 166–169 reinforced elastomeric materials, 333 thermoplastic polyurethane nanocomposites, 489–491 dispersion, 501–503 Window channels (automotive), rubber applications, 529 X-ray diffraction analysis: apolar diene rubber nanocomposites, dispersion characterization, 383–387 ethylene-propylene-diene-monomer rubber, in situ intercalation preparation, 469–473

601

rubber-clay nanocomposite morphology, 182 Zero-shear viscosity, polyepichlorohydrin rubber-clay nanocomposite rheology, 259–261 Zinc dimethyl dithiocarbamate (ZDMC), vulcanization and permeability, 356–358 Zinc oxides: melt compounding, rubber-clay nanocomposites, 165–169 rubber-clay nanocomposites, vulcanization activation, 138–141 vulcanization reaction, organo-clay kinetics, 285–289